Silicon heterojunction solar cells and methods of manufacture

ABSTRACT

The present invention relates to a solar cell comprising a heterojunction photoelectric device comprising, a front electrode layer, a back electrode layer comprising a metallic contact layer, a light-absorbing silicon layer arranged between said front electrode and said back electrode layers and a doped silicon-based layer arranged between said light-absorbing silicon layer and said back electrode layer, characterized in that said heterojunction photoelectric device further comprises a wide band gap material layer having an electronic band gap greater than 1.4 eV, said wide band gap material layer being applied on a surface of the light-absorbing silicon layer between said light-absorbing silicon layer and said doped silicon-based layer. The present heterojunction layer or stack of layers is compatible with thermal annealing and firing processes at T above 600° C.

TECHNICAL FIELD

The present invention relates to the field of solar cells. Moreparticularly, the present invention relates to a photovoltaic solar cellcomprising a heterojunction silicon based photovoltaic device comprisinga heterojunction obtained from a low cost metallization fabricationmethod.

PRIOR ART

Today's solar cell production is dominated by two major silicon solarcell structures. These are aluminium-back surface field (Al-BSF) solarcells and the quickly emerging passivated emitter and rear solar cells(PERC), together accounting for more than 90% of the world production in2016. PERC cells, represented in FIG. 1A, are very similar to Al-BSFcells, but have improved performance due to a dielectric layer on therear side of the cell, i.e. on the side of the cell farthest from thatbeing most irradiated by impinging light. The dielectric layerpassivates the silicon wafer surface and acts as internal mirror. Thefabrication of PERC solar cells starts with boron B-doped wafers(p-type) that are wet-chemically textured and cleaned, and usuallyrelies on at least one high-temperature step. The at least one hightemperature step comprises the diffusion of the n-type front contact inan atmosphere of POCl₃ at temperatures of ca. 850° C. This stepmassively improves the lifetime of minority charge carriers in manysilicon materials, for example in multi-crystalline silicon material,which is linked to gettering of impurities. For the avoidance of doubts,the terms “multicrystalline material” shall be understood here asreferring to any material consisting of more than one crystal grain, incase of the material being a silicon wafer to multicrystalline,polycrystalline, quasi-monocrystalline, and other type of siliconwafers. After applying a silicon nitride (SiN_(x)) anti-reflectioncoating to the front, the rear side is cleaned and coated with adielectric layer (for example AlOx/SiNx stack) for surface passivation.Passivation of the silicon wafer or at least a surface portion thereofis acknowledged in the field of solar cell design and manufacturing toimprove the minority charge carrier lifetime of the siliconwafer.Because non-conductive the rear passivation layer needs to bepatterned. After this step, front- and back-metallization is applied byscreen-printing of silver- and aluminium-pastes, respectively.Subsequently, a second high-temperature step, also referred to as“co-firing”, is employed. During this second high temperature step theglass frit of the front metallization first sinters through the SiN_(x)anti-reflection coating to establish the front contact. Secondly,passivating species, for example any of hydrogen, deuterium, tritium,neutral or charged is released from the SiN_(x) layer and partlydiffuses into the silicon wafer where it passivates broken silicon bondsand other defects, which is highly beneficial in case ofmulticrystalline wafers but also for monocrystalline material. Thefiring step requires rapid ramping (and cooling) to a peak temperature,typically a few seconds at 750° C.≤T≤850° C. This co-firing step iscritical for contact formation and has massively contributed to increasethroughput and decrease cost of industrial wafer based c-Si solar cells(Al-BSF and PERC), making this technology mainstream in the photovoltaicmarket.

However, the performance of PERC solar cells is inherently limited bythe two dimensional charge carrier flow defined by the rear geometry andby charge carrier recombination at the direct metal-silicon contacts(black arrows in FIG. 1A) [Cuevas, A. Physical model of backline-contact front-junction solar cells. 164502, (2014),DOI:10.1063/1.4800840].

In fact, surface recombination velocity is extremely high (>10⁵ cm/s) atsuch direct metal-Si contacts, leading to a drop in the minority chargecarriers quasi-Fermi levels and consequently to a voltage loss, whichcan be higher than 60 mV. In order to avoid such voltage (and thusefficiency) loss, the interface between the absorber (Si wafer) and themetal has to be electrically passivated (FIG. 1B). Thin buffer layerssuppress the recombination of minority charge carriers by displacing themetal contact from the Si wafer. At the same time, the buffer layer issufficiently conductive to extract the majority charge carriers, whichare selectively contacted by a doped layer on top.

A prime example of passivated contacts is found in a-Si:H/c-Siheterojunction (SHJ) solar cells, pioneered by Sanyo/Panasonic over thepast 2 decades (“Heterojunction with Intrinsic Thin layer”, HIT™), andcontinuously optimized until reaching efficiencies up to 24.7% in 2013.The wafer surface passivation of silicon heterojunction solar cellsrelies on saturation of dangling bonds by amorphous silicon, morespecifically by hydrogen contained in the amorphous silicon layers. Thehydrogen content is typically between 8% and 40% [R. A. Street,Hydrogenated Amorphous Silicon, Cambridge University Press, 1991,https://doi.org/10.1017/CB09780511525247]. The large amount of hydrogenresults in lower mass density compared to de-hydrogenated silicon, andalso a decreased refractive index.

Even though impressive results have been obtained with SHJ solar cells,they feature several major drawbacks.

Upon heating a-Si:H layers, hydrogen effuses. As hydrogen dissolvessuddenly from the Si atoms of the a-Si:H layer, diffusion through thesilicon layer can limit the effusion, which can lead to the formation ofblisters.

Additionally, hydrogen effusion becomes an irreversible process when theSHJ layers are annealed at temperature higher than 350° C., andintroducing hydrogen after deposition or re-introducing hydrogen aftereffusion cannot restore an initial good electrical quality, but leads tobad electrical quality [J. Shi, Appl. Phys. Lett. 109, 153503153503(2016); 10.1063/1.4964835].

As already mentioned above, cost-effective industrial metallizationschemes are mainly based on firing-through of high temperature Ag-pastesor by plating of Ni/Cu. In both cases process temperature are well above250° C. and thus incompatible with SHJ. In more details, Ag-basedmetallization require a firing through step, few seconds at 750°C.≤T≤850° C. While Ni/Cu metallization then requires a sintering step ataround 350° C. to form a silicide. In addition, surface passivation byamorphous silicon further requires locally very flat surfaces. Usuallysolar cell surfaces are textured with pyramids to enhance lighttrapping, for example random pyramids, the edges of the pyramids need tobe rounded and their sidewalls polished in order to smoothen the surfaceto attain good surface passivation.

Another drawback of this passivating contact technology arises fromparasitic light absorption in the a-Si:H films and transparentconductive oxides (TCOs), which induces current losses of about 3mA/cm². These parasitic absorption effects cannot be reduced any more byfurther decreasing the film thickness because minimal layer thickness isneeded to ensure sufficient later conductivity.

Hence, production of SHJ cells requires investing in a completely newmanufacturing line and interconnection technology to the low temperaturemetallization that is not proven in the field for decades. This leavesout to cell manufactures the option to upgrade existing tools. Withinthe uncertain financial times for the photovoltaic industry, thistranslates into a too high investment risk representing the main hurdlefor adopting of this cell technology.

An attractive industrial solution would be to combine the highlytransparent front side structure of a PERC or Al-BSF solar cell with anovel rear passivating contact with passivation properties as good asSHJ and that are stable to high temperature thermal treatments used forfabrication of industrial metallization schemes. Additionally, whenreducing recombination rate of the rear passivating contact, front siderecombination at the homo-junction contact becomes dominant. To mitigatesuch losses, highly-transparent and temperature stable passivatingcontacts for front side are required. In this case, the high temperaturetreatment would also increase layer crystallinity thus reduce opticalparasitic absorption compared to SHJ. This passivating contacttechnology, thus, will offer to solar cell manufacturers the option ofintroducing high efficiency passivating contact concepts on at least oneside of the solar cell without the need of radically change theirproduction lines.

DISCLOSURE OF THE INVENTION

According to a first aspect the present invention relates to a solarcell comprising a heterojunction photoelectric device comprising, afirst electrode layer, a second electrode layer comprising a metalliccontact layer, a light-absorbing silicon layer arranged between saidfirst electrode and said second electrode layers and a passivating andconductive stack comprising at least one doped silicon-based layerarranged between said light-absorbing silicon layer and at least one ofthe first and second electrode layer, characterized in that thepassivating and conductive stack further comprises at least one wideband gap material layer having an electronic band gap greater than 1.4eV, said wide band gap material layer being applied on a surface of thelight-absorbing silicon layer between said light-absorbing silicon layerand said doped silicon-based layer, thereby forming a heat-resistantheterojunction contact, said heat resistant heterojunction beingarranged for at least maintaining its passivating and conductiveproperties after thermal treatment thereof above 600° C.

In the context present invention, the terms “conductive stack” shall beconstrued as a layer stack in a heterojunction solar cell havingspecific contact resistivity ranging from 1 and 10⁴ m ohm·Cm².

The solar cell of the invention therefore comprises a heat-resistantheterojunction contact at any of the rear side, the front side, or boththe rear side and the front side of the cell, or two heterojunctioncontacts on the same side, preferably the rear side, of the solar cell,i.e. between the light absorbing silicon layer and the metallizationlayer, offering greater surface passivation than known solar cells,boosting the open-circuit voltage (Voc) potential from 680 mV to 740 mV,while PERC cells are limited to 660 mV and up to 680 mV.

In the context of the present invention the light absorbing siliconlayer can be made of n-type or p-type or intrinsic silicon withresistivity comprised between 0.1 and 1000 ohm-cm.

In the present invention, the terms “heat-resistant heterojunctioncontact” shall be understood as any heterojunction contact in a solarcell structure capable to at least maintain or improve its passivatingand conductivity properties when exposed to thermal treatments.

Thermal treatments in the context of the present invention shallcomprise as “firing” thermal treatment when based on heating and/orcooling ramps above 20° C./s at temperatures of 600° C. to 1000° C. anddwell time of 1 s to 10 min. Thermal treatments shall also comprise“thermal annealing” when based on heating and/or cooling ramps between 1and 100° C./min, at temperatures of 600° C. to 1000° C. and dwell timeof 1 min to 60 min.

In a first embodiment of the inventive solar cell, an interface betweensaid passivating and conductive stack and the light absorbing siliconlayer has a very abrupt or shallow distribution of doping profile, saidpassivating and conductive stack and thermal treatment being configuredsuch that to prevent dopants to diffuse from the passivating andconductive stack to the light absorbing silicon.

The terms “very abrupt or so-called shallow distribution” shall beunderstood in the context of the invention as the intensity [counts/s]of the dopants into the light absorbing silicon decays over at least oneorders of magnitude within a distance of less than 20 nm, preferablyless than 10 nm, further preferably less than 5 nm.

In a second embodiment the light absorbing silicon layer comprises adoped region with sheet resistance between 1 and 10⁵ ohm/square at aninterface between said passivating and conductive stack and the lightabsorbing silicon layer, said doped region being obtained by diffusionof dopants from any of the layers forming the passivating and conductivestack into to the light absorbing silicon layer during said thermaltreatment.

In a preferred embodiment, the heat resistant heterojunction is appliedon a rear side of the solar cell.

In preferred embodiments, the wide band gap material layer has athickness of at most 20 nm, preferably between 0.5 nm and 2 nm.

In various embodiments, the wide band gap material layer comprises adielectric material, which is advantageously chosen among any of thefollowing materials: SiOx, SiCx, SiNx, SiOxNy, SiOxCy, SiCxNy, SiOxCyNz,AlOx, HfOx, AlHfOx, AlNx, TiNx, ZrOx, Y2Ox, AlSiOx, HfSiOx, AlHfSiOx,where x, y, and z are entire numbers, and which can be amorphous orcrystalline.

Advantageously, any Si-based materials containing O, N, or C in anycomposition forming the wide band gap material layer may also containfurther elements, and can be amorphous or crystalline.

In preferred embodiments, the passivating and conductive stack has aneffective refractive index lower than that of amorphous silicon at awavelength 633 nm. For the avoidance of doubts, the terms “effectiverefractive index” means the refractive index of the passivating andconductive stack determined according to an effective mediumapproximation method.

Further, any of the previously recited wide band gap preferred materialsmay contain hydrogen, fluorine, phosphorous, boron, and other elements,and preferably containing at least one of hydrogen or fluorine.

In a preferred embodiment of the invention, the wide band gap materiallayer is made of SiOx, where x is being chosen between 0.5 and 2.

In another preferred embodiment of the invention the doped silicon-basedlayer has an atomic percent concentration of hydrogen of less than 5% asthe number of hydrogen atoms per unit volume divided by the total numberof all atoms per unit volume of the doped silicon-based layer, i.e.meaning a hydrogen concentration, which is more than one order ofmagnitude lower than the hydrogen concentration of a-Si:H, which istypically in the range 8%-40% [Beyer 2016, 10.1002/pssa.201532976]. Moregenerally speaking, it is preferred that the doped Si-based layer or apart of it is being more transmissive than silicon itself for molecularor atomic hydrogen.

In other embodiments the passivating and conductive stack of theinvention further comprises a buffer layer arranged between said wideband gap material layer and said doped silicon-based layer, which bufferlayer is preferably itself also a silicon-based layer and can be dopedor intrinsic.

Various materials can be chosen to realize the buffer layer but it ispreferred according to the invention the said buffer layer is made of atleast one of the materials chosen among Si, SiCx, SiNx, SiOx, SiCxNy,SiCxOy, SiNxOy, SiCxNyOz or a combination thereof.

The inventors of the current invention have identified thatappropriately chosen material for the wide band gap layer canadvantageously provide passivation of the light-absorbing siliconsurface without requiring further re-hydrogenation of thelight-absorbing silicon surface and or wide band gap layer after thermaltreatment. This holds for the case with and also without buffer layer,directly or indirectly between the doped-silicon based layer and thelight absorbing silicon.

In addition, where said light-absorbing silicon surface is notpassivated enough after thermal treatment the passivating and conductivestack and/or the capping layer may comprise(s) passivating species topassivate defects within the said passivating and conductive stackand/or at the interface between said passivating and conductive stackand said light absorbing silicon layer, said passivating species beingreleasable upon said thermal treatment.

In particular, the inventors have found out that addition of fluorine tothe layer stack can enhance passivation even without furtherhydrogenation. Such fluorine can be added already in the growth of thelayer stack during manufacturing of the solar cell structure, forexample using a fluorine compound as precursor gas or target, orafterwards for example by implantation.

Adding fluorine already in the growth of the layer stack is especiallyadvantageous because this reduces the total number of process steps. Inaddition, during thermal treatment, the fluorine can be redistributedand can diffuse to the interface between the light absorbing siliconlayer and the wide band gap material layer, enhancing passivation.

Hydrogen can also have other effects: Molecular or atomic hydrogen canbond to dopant impurities, deactivating the dopant, and hydrogen canbond to fixed charges and influence the fixed charge density.

Importantly, both effects can influence the band bending in thelight-absorbing silicon. According to the invention hydrogenation isutilized to influence doping, fixed charges, or band bending.

According to another characteristic of the invention a transparentconductive oxide layer may be arranged between said doped silicon-basedlayer and said metallic layer.

In another preferred embodiment of the solar cell of the invention thesurface of the light absorbing silicon layer whereupon the wide band gapmaterial layer is applied may be structured, i.e. showing a determinedand controlled surface pattern or structure providing determined textureand/or roughness of the surface. Any kind of surface texture and/orroughness of the light absorbing silicon layer may be contemplated inthe context of the invention, in particular nanoscale texturing tocreate surface roughness at atomic level.

In other preferable embodiments the wide band gap material layer maycomprise through-holes extending from said first surface to said secondsurface.

Additionally, in any embodiment of the invention the solar cell may alsocomprise a capping layer arranged between said doped silicon-based layerand said metallic contact layer of the second electrode.

In alternative embodiments of the present invention, the wide band gaplayer, the buffer layer or the silicon-based doped layer may be comprisea so-called heterogeneous layer as described in co-pending PCTapplication WO 2017182472, the content of which is herewith fullyincorporated by reference.

In a further embodiment the passivating and conductive stack and/or thecapping layer may comprise(s) passivating species as defined above topassivate defects within the said passivating and conductive stackand/or at the interface between said passivating and conductive stackand said light absorbing silicon layer, said passivating species beingreleasable upon said thermal treatment.

According to a second object, the present invention also relates to amethod for manufacturing a solar cell as previously presented, saidmethod comprising a step of depositing a passivating and conductivestack comprising a wide band gap material layer on a surface of a lightabsorbing silicon layer, a step of depositing a doped silicon basedlayer on said wide band gap material layer and then at least one thermaltreatment of the stack formed thereby at above 600° C. in order toreleases passivating species from the said passivating and conductivestack and/or at an interface between said passivation stack and saidlight absorbing silicon layer. Additionally, the thermal treatmentincreases crystallinity of the silicon-based doped layer enhancing itsoptical and electrical properties.

Preferably, the wide band gap material layer has an electronic band gapgreater than 1.4 eV before and after thermal treatment.

Preferably, the electronic band gap of the wide band gap material layervaries less than 20% during or after said thermal treatments.

Preferably, the thermal treatment comprises a firing step of the solarcell at temperatures above 600° C. and with at temperature ramp rates ofmore than 20K/s and dwell time at maximum temperature below 10 s (rapidthermal annealing).

Such thermal treatment thereby allows for reducing hydrogen content ofthe doped silicon-based layer because the elevated temperatures causethe rupture of hydrogen bonds. In a preferred embodiment, thecombination of thermal treatment and the passivating and conductivelayer stack prevent in-diffusion of dopant impurities form thesilicon-based layer through the wide band gap layer to thelight-absorbing silicon layer.

In preferred embodiments, either during the first treatment or anotherthermal treatment, the hydrogen content of the wide band gap materiallayer or the light absorbing silicon layer interface is increased withrespect to the state of the respective layer before that thermaltreatment. Thermal treatment can result in a gradual decrease of thelayer stack crystallinity when moving towards the side of the layerstack facing away from the light absorbing silicon.

Finally, the method of the invention may further comprise the steps of:

-   -   forming a capping layer between said doped silicon-based layer        of the passivating and conductive stack and said metallic        contact layer; and    -   thermally treating the passivating and conductive stack and the        capping layer in order to release passivating species from any        of the passivating and conductive stack and the capping layer to        passivate defects within the said passivating and conductive        stack and/or at the interface between said passivating and        conductive stack and said light absorbing silicon layer.

PRESENTATION OF DRAWINGS

Various embodiments of the present invention will now be described inrelation to the appended drawings, among which:

FIGS. 1A and 1B represent various solar cell structures known from theprior art,

FIGS. 2 A and B represents first embodiments of a solar cell accordingto the invention,

FIG. 3A to 3C represent the influence of a capping layer to controltransmission of a passivation component such as hydrogen or fluorine anddensity of fixed charges within the passivating and conductive stackstructure of a solar cell according to the invention;

FIG. 4A to 4D represent various alternative constructions of apassivating and conductive stack for a solar cell according to theinvention in a second embodiment;

FIGS. 5A and 5B represent a third embodiment of a passivating andconductive stack for a solar cell structure according to the invention;

FIG. 6A to 6D represent various alternatives of a passivating andconductive stack for a solar cell structure according to a fourthembodiment of the invention;

FIG. 7A to 7D represent other embodiments of a solar cell structureaccording to the invention, where the solar cell is a rear sidecontacted solar cell;

FIGS. 8A and 8B represent the dependence of the value of effectiveminority charge carrier lifetime values of symmetric sister samples on(a) the NH₃/SiH₄ flow ratio during the SiNx deposition, and (b) the SiNxthickness. “Sister samples” means that samples were co-processed in allother process steps;

FIGS. 9A and 9B represent, in the first example of the contact structureof the solar cell of the invention, thermal desorption spectroscopy(TDS) spectra of as deposited a-SiCx (p) on SiOx/c-Si, SiNx on c-Si, andSiNx on SiCx(p)/SiOx/c-Si, with the SiCx(p) previously annealed at 800°C., respectively for m/Z=2 (H₂), and mass separated ions (m/z) m/Z=18(H₂O), obtained by heating the sample at a ramping rate of 20° C./minafter a vacuum of 5.25×10⁻¹⁰ Torr was reached in the measurementchamber;

FIGS. 10A to 10C represent, for the first example of the contactstructure of the solar cell of the invention, implied open circuitvoltage (iVoc) values for samples produced by varying the trimethylboron(TMB) flow during the doped layer deposition from 0.2 standard cubiccentimeters per minute (sccm) to 1.9 sccm and annealed at temperaturesrange of 775° C. to 900° C. with various annealing dwell times of 0minute, 5 minutes, and 15 minutes and using heating and cooling ramps of10 and 2° C./min, respectively;

FIGS. 10D to 10F represent for the first example of the contactstructure of the solar cell of the invention measured specific contactresistivity values as a function of TMB flow used during the depositionof the SiCx(p) for different thermal annealing dwell temperatures anddwell times of 0 minute, 5 minutes, and 15 minutes;

FIG. 11 show the doping profile within the c-Si wafer in of the firstexample of the solar cell of the invention measured by ElectrochemicalCapacitance-Voltage (ECV) for different thermal annealing dwelltemperatures and dwell times for the contact that was produced with theTMB flow of 1.5 sccm on top of chemical SiOx and 10 nm thick intrinsicSi interlayer;

FIG. 12A-12D represents output characteristics of the hybrid solar cellsin the first example of the invention exhibiting the boron doped layerproduced with the TMB flow of either 1.5 or 1.9 sccm and thermal annealat dwell temperatures of 800° C., 825° C., and 850° C.;

FIGS. 13 to 20 represent various configurations and measurement resultsfor a wide-band gap silicon carbide for front side carrier selectivecontacts as a second example of the invention:

FIG. 21 represent the solar cell hybrid structure, TEM image of theSi/SiCx(p) layer stack after thermal annealing and external parametersof the cell;

FIG. 22A-D report the conductive atomic force microscope (c-AFM) of theSi/SiCx(p) layer stack after thermal annealing at 850° C. and differentTMB flows;

FIG. 23A reports the c-AFM of the Si/SiCx(p) with TMB flow of 1.5 sccmsin the as-deposited state; B, C and D after thermal annealing of 800 C,850 and 925° C.;

FIG. 24A-24B. TEM and EDX maps of the Si/SiCx(p) layer stack afterannealing at 850 and 925° C.

FIG. 25A shows the implied open circuit voltage as a function of processgasses flux ratio (CH₄/(SiH₄+CH₄+BF3)) and effect of fluorine aspassivating agent in a solar cell according to the invention;

FIG. 25B show the related specific contact resistivities for theselected samples from the same sample set.

FIG. 26 shows the lifetime evolution of the sample produced withCH4/(SiH4+CH4+BF3) ratio of 0.1 and thermal annealing at differenttemperatures from 800° C. to 875°;

FIG. 27 reveals the scanning transmission electron microscopy imageswith Energy-dispersive X-ray spectroscopy (EDX) mapping for Si, O and Fatoms for SiCx doped with BF3 in the as-deposited state, after thermalanneal at 850 and 925° C. for 5 min;

FIG. 28A represents cell parameters obtained using fluorinated dopantsource (BF₃) for hole selective contact fabricated with a thermal annealwith and without post-hydrogenation;

FIG. 28B represents the cell parameters for the best cell;

FIG. 29 represents the process flow for the fabrication of thesymmetrical structure used for the development of the firedheterojunction in the fourth example of this invention;

FIG. 30 shows minority carrier lifetime (τ_(eff)) at injection level of10¹⁵ cm⁻³ as function of the normalized flow ratio between CH4 and totalgas flow (i.e. CH4+SiH4+TMB+H2) of the SiCx(p) layer forming the firedheterojunction;

FIG. 31 minority carrier lifetime (τ_(eff)) as function of the minoritycarrier density for the optimal layer conditions (normalized flow ratiobetween CH4 and total gas flow of 0.18 for the SiC_(x) layer) with andwithout a-Si(i) buffer layer;

FIGS. 32A-32C represent Secondary-ion mass spectrometry (SIMS)measurements of the O, B and C doping profiles expressed in intensity[counts/seconds] as function of the depth, for the SiCx layer formingthe passivating and conductive stack of the invention with normalizedflow ratio between CH4 and total gas flow of 0.18 in the as-depositedstate and after firing step;

FIGS. 32D-32F represent SIMS of the O, B and C doping profiles expressedin intensity [counts/seconds] as function of the depth, for thepassivating and conductive stack of the present invention contact formedby the intrinsic buffer layer capped with the SiCx with normalized flowratio between CH4 and total gas flow of 0.18 in the as-deposited stateand after firing step;

FIG. 33A represents scanning transmission electron microscopy (STEM)high-angle annular dark-field imaging (HAADF) image of the c-Si withSiOx/nc-SiCx(p) fired heterojunction, shown alongside an EDX map and aline scan of the Si K and O K and C K edge EDX intensities;

FIG. 33B represents STEM HAADF image and EDX map of the c-Si withSiOx/nc-Si(i)/nc-SiCx(p) double layer forming the fired heterojunction,also shown alongside the corresponding EDX line scan of the Si K and O Kand C K intensities;

FIG. 33C represents high resolution transmission electron microscopy(HRTEM) micrograph of the c-Si with SiOx/nc-SiCx(p) fired heterojunctionand corresponding Fourier transforms computed in the first 7 nm of thenc-SiCx(p) close to the SiOx (inset i) and in the upper part of thedoped layer (inset ii);

FIG. 33D represents HRTEM micrograph the c-Si withSiOx/nc-Si(i)/nc-SiCx(p) double layer fired heterojunction;

FIG. 34 represents Fourier-transform infrared spectroscopy (FTIR)measurements of the of the fired heterojunction formed by asilicon-based doped layer of SiC_(x)(p) with normalized flow ratio of0.18% in the as-deposited, fired and fired followed by hydrogenation;

FIG. 35 represents TDS effusion profile of the (m/z=2) of the firedheterojunction formed by a silicon-based doped layer of SiC_(x)(p) withnormalized flow ratio of 0.18 in the as-deposited, fired and fired,followed by hydrogenation;

FIG. 36 represents light soaking under 1-sun illumination for the firedheterojunction employing as silicon-doped layer a SiCx(p) withnormalized flow ratio of 0.78, 0.18 with and without buffer layer;

FIGS. 37A and 37B represents sketches of two hybrid solar cellstructures involving a-Si(i):H/a-Si(n):H as front side emitter and thepresented fired heterojunction with the silicon-doped layer the SiCx(p)or its variation with an a-Si(i)(buffer)/SiCx(p) on the rear siderespectively with normalized flow ratio of the SiCx(p) of 0.18;

FIG. 37C represents current density-voltage (J-V) measurements ofproof-of concept hybrid cells involving a silicon heterojunction on thefront side and fired heterojunction with the silicon-doped layer theSiCx(B) or its variation with an a-Si(i)(buffer)/SiCx(p) on the rearside respectively with normalized flow ratio of the SiCx(p) of 0.18;

FIG. 38 represents light soaking after six consecutive measurements ofthe cells employing a conventional silicon heterojunction on the frontside and the fired heat resistant heterojunction on the rear side Asingle nc-SiCx(p) with normalized flow ratio of 0.18 fired at twodifferent peak temperatures for a dwell time of 3 sec, B singlenc-SiCx(p) as function of the normalized to the max value of the TMB_2%flow; C single nc-SiCx(p) with normalized flow ratio of 0.18 fired atsame peak temperatures and different dwell time, D bilayernc-Si(i)/nc-SiC(p) with normalized flow ratio of of the SiC 0.18 asfunction of the normalized to the max deposition time of the intrinsicbuffer layer.

FIG. 39 represents Fourier-transform infrared spectroscopy (FTIR)measurements of the fired heterojunction with employing SiC_(x)(p)layers with different CH₄ flows with and without aSi(i) as buffer layerafter firing and hydrogenation. Si—C and C—H_(n) absorption peaksincrease with increasing of the normalized flow ratio CH₄ flow;

FIG. 40 represents HRTEM micrograph, STEM HAADF image and EDX map of thec-Si with SiOx/Si/SiCx(p) after thermal annealing;

FIG. 41 represents i-VOC and J0 as function of the thickness of the thea-Si buffer layer forming the a-Si(i)/a-SiCx(p) (in the as-depositedstate) after thermal annealing;

FIG. 42 represents surface photo voltage (SPV) measurements of the firedheat resistant heterojunction with single SiCx(p) with several values ofthe normalized flow ratio before and after in-situ light soaking; SPV ofthe fired heat resistant heterojunction with single SiCx(p) withdifferent values of the normalized TMB flow ratio before and afterin-situ light soaking; SPV of the fired heat resistant heterojunctionwith bilayer Si(i)/SiCx(p) with different values of the normalizeddeposition time for the Si(i) buffer layer before and after in-situlight soaking.

FIG. 43 Represents external parameters (i.e. Jsc, VOC, FF and η) ofhybrid cells with different firing temperature of the heat resistantheterojunction;

FIG. 44 represents PLI measurements of hybrid cells with differentfiring temperature of the heat resistant heterojunction before and aftercuring of the ITO (i.e. annealing at 210° C. for 30 min);

FIG. 45 represents external parameters of hybrid cells with differentdoping flow of the heat resistant heterojunction;

FIG. 46 represents PLI measurements of hybrid cells with differentdoping slow of the heat resistant heterojunction before and after curingof the ITO (i.e. annealing at 210° C. for 30 min);

FIG. 47 represents external parameters of hybrid cells with differentfiring time of the heat resistant heterojunction;

FIG. 48 represents PLI measurements of hybrid cells with differentfiring time of the heat resistant heterojunction before and after curingof the ITO;

FIG. 49 represents external parameters of hybrid cells with differentthickness of the intrinsic buffer layer of the heat resistantheterojunction;

FIG. 50 represents PLI measurements of hybrid cells with different withdifferent thickness of the intrinsic buffer layer of the heat resistantheterojunction before and after curing of the ITO;

FIG. 51 represents alternative hydrogenation process of the heatresistant heterojunction

FIG. 52 represents a process flow for manufacturing p-type PERC (steps1-6) compared with the inventive FlaSH contact integration (steps1-6.a).

DESCRIPTION OF PREFERRED EMBODIMENTS

The present invention relates to the provision of improved siliconheterojunction solar cells and methods for the manufacturing thereof.

As will be understood from the following description and the appendedfigures the solar cell of the invention relies on the provision, in aheterojunction-type solar cell, of a passivating and conductive stackcomprising a thin wide band gap material layer, formed of a materialsuch as SiOx for example, as a passivation layer between a lightabsorbing silicon layer, preferably formed of a silicon wafer and adoped silicon-based layer.

FIGS. 2 to 7 represent a cross section of a heat-resistantheterojunction solar cell HC according to the invention in variousembodiments. The heterojunction solar cell HC essentially comprises alight absorbing silicon layer 1, for example consisting of a siliconwafer, a first and a second electrode layers 300, 301 and a passivatingand conductive stack PC arranged between said light absorbing siliconlayer 1 and at least one of the electrode layers. According to theinvention, the passivating and conductive stack PC comprises at least awide band gap material layer 10 arranged on a surface of the lightabsorbing silicon layer 1 and doped silicon-based layer 6 arranged on asurface of the wide band gap layer 10. In the context of the invention,the wide band gap material layer 10 is defined as a material having anelectronic band gap greater than 1.4 eV, the nature and function ofwhich will be described hereinafter. Optionally, the passivating andconductive stack PC may also comprise a buffer layer 4 and a cappinglayer 12.

The use of such thin wide band gap material layer 10 in combination witha doped silicon-based layer 6, preferably comprising silicon carbideSiCx, as passivating contact instead of intrinsic hydrogenated amorphoussilicon commonly found in standard silicon heterojunction solar cellsprovides a more homogenous and more temperature stable passivatingcontact allowing for the heterojunction solar cells of the invention towithstand firing and/or thermal annealing (as defined above) making theinventive solar cell and heterojunction compatible with low costmetallization schemes required for industrialization of heterojunctionsolar cells.

As opposed to hydrogenated amorphous silicon, the wide band gap materiallayer 10 avoids epitaxial growth of the doped Si-based layer 6 andallows to achieve high surface passivation after and/or during firingand/or thermal annealing. For example, upon firing of the heterojunctionlayer stack of the invention a passivating species (i.e. hydrogen,fluorine, nitrogen, or oxygen) contained in doped silicon based layer 6or another layer forming the passivating and conductive stack, evolveand incorporate in the wide band gap material layer 10 or at theinterface between the light absorbing silicon layer 1 and wide band gaplayer 10 and reduces the interface state density.

According to the invention, surface passivation of the proposedheterojunction is activated by the thermal treatment process. In case offiring, this allows to make use of co-firing processes which areemployed in manufacturing of most industrial wafer based c-Si solarcells to form the front contact and the high quality passivation rearpassivating and conductive stack (heat-resistant heterojunction).Additionally, the HC stack ensures full area contact functionalities(i.e. no need for patterning), thus reducing processing steps whileincreasing conversion efficiency with respect to PERC solar cells.

Passivating species preferably comprise according to the presentinvention chemical elements such as hydrogen, fluorine, nitrogen, oroxygen. Such passivating species may provide passivation of a surface,as well as of electronic defects, through interfacing between layersand/or chemical reactions.

In the present invention, firing can also be used to hydrogenate thewide band gap material layer 10 and decrease its defect density and theinterface with the light absorbing silicon layer 1. For a SiOx wide bandgap material layer 10, the optimum temperature for hydrogenincorporation at the Si-SiOx interface is higher than 400° C. comparedto 200° C. of a-Si:H, which illustrates the enhanced temperaturestability compared to silicon heterojunctions known from the prior art.Hydrogen incorporation can also be achieved by rapid thermal annealing(firing) at temperatures higher than than 600° C., still preferablyhigher than 750° C. by first providing a hydrogen rich wide band gaplayer 10, for example SiNx, and then firing. As demonstrated by theinventors hydrogen can be also incorporated during the firing of theheat-resistant heterojunction.

Depending on the density of defects at the interface electrons or holesinjected from the silicon wafer 1 to the wide band gap material layer 10during device operation or irradiation can be trapped as fixed chargesor deactivate fixed charges, resulting in more favourable alignment ofthe energy bands. This means that the height of the potential barrierfor the collected carrier decreases, thus enhancing surface passivationand charge carrier transport.

The charges can be trapped or deactivated at the light absorbing Siwafer 1-wide band gap material layer 10 interface, in the wide band gapmaterial layer 10, at the wide band gap layer-doped silicon based layer6 interface, or where the passivating and conductive stack PC comprisesa buffer layer 4 and/or capping layer 12 in the wide band gap materiallayer 10-buffer layer 4 interface, or in the buffer layer 4 or in thedoped silicon based layer 6, or in the capping layer 12, when a cappinglayer 12 is provided in the inventive stack as shown in FIGS. 3 and 6 inparticular.

As opposed to amorphous silicon interface layers prepared by PECVD andalready employed in some SHJ solar cells the wide band gap layer 10,which may be made of SiOx in a preferred embodiment, of the inventiveheterojunction does not necessarily passivate the light-absorbingsilicon layer 1 surface directly upon application thereon, but onlyafter thermal treatment(s) of the passivating and conductive stack ofthe heterojunction. The capacity of the inventive heterojunction solarcell HC to undergo and withstand thermal treatment(s) and to provideincreased electrical generation efficiency thereafter is thus a centralaspect of the of an heterojunction solar cell according to the inventionand its manufacturing process.

Such SiOx wide band gap layer 10 furthermore has the advantage of beingmore transparent than a-Si:H layers, which is advantageous for frontside application of the inventive heterojunction.

The wide band gap material layer 10 prevents epitaxial regrowth of thedeposited layer during high temperature (T>600° C.) thermal treatment.It is preferably provided as a thin wide band gap material layer 10,i.e. having a thickness of at most 20 nm. When said wide band gap layer10 is realized from SiOx material, its thickness is usually less than 3nm, preferably less than 1.5 nm, which allows charge carrier transportthrough the layer.

Optionally, pinholes are formed in the thin wide band gap material layer10 during the process, leading to a direct connection between thelight-absorbing silicon layer 1 and the doped silicon-based layer 6and/or any further deposited silicon-based buffer layer 4. The size ofthe pinholes is typically in the range of the thickness of thedielectric layer, or in the range 0.2 nm to 200 nm, preferably 0.2 nm to20 nm.

With the heterojunction solar cell HC and firing thermal treatment ofthe present invention, dopants may not diffuse from the dopedsilicon-based layer 6 through the wide-band gap layer 10 into thelight-absorbing silicon layer 1. In the latter case, a very abrupt,so-called shallow, buried junction is formed i.e. intensity [counts/s]of the dopants decaying over at least one orders of magnitude within adistance of less than 20 nm, preferably less than 10 nm, furtherpreferably less than 5 nm. The advantage of such very abrupt shallowburied junctions is that the total dopant dose in the silicon wafer 1 ismuch lower than in highly-doped regions employed commonly, whichtypically have a depth of above 20 nm with a surface dopingconcentration of more than 5.10¹⁹ cm³, because the lower dopant dosereduces, or eventually eliminates completely, Auger recombination andfree-carrier absorption.

Preferably, the wide band gap layer 10, the silicon-based buffer layeror the silicon-based doped layer 6 can be doped and act as doping sourcefor the wafer, i.e. the light-absorbing silicon layer 1 during a thermaltreatment.

This effect relies on the following physical phenomenon. A segregationcoefficient k for the dopant can be ≠1 for the materials in the (A)light-absorbing silicon layer 1 or one of the other layers, and (B) oneof the at other layers.

k=C _(eq,A) /C _(eq,B)

C_(eq,A) is the equilibrium concentration of the dopant in material A,and

C_(eq,B) is the equilibrium concentration of the dopant in material B.

For k>1, dopants can be expelled from B to A during a thermal treatment,even if before said thermal they are present in equal concentration isboth A and B, resulting in a step in the doping profile. See also [Sze,Semiconductor Devices: Physics and Technology, John Wiley & Sons Ltd,2012, ISBN 978-0470537947] for a more detailed discussion. This effectcan be utilized for instance to expel dopants during a thermal treatmentfrom one of the layers of the layer stack to the light-absorbing siliconlayer 1, potentially useful to increase doping concentration in thelatter. Consequently, the doping profile in the light-absorbing siliconlayer 1 or in the passivation layer stack (formed of all layers on saidsilicon layer 1) can depend on the properties (composition, doping,thickness, etc.) of the silicon-based doped layer 6 or silicon-basedbuffer layer 4, also independent of the thermal treatment. Thepassivating and conductive stack PC can thus be tuned towards achievinga specific doping profile in the wafer 1 or some other adjacent layerwithout changing the conditions of the thermal treatments.

The wide band gap layer 10 is usually a thin passivating layer, in someembodiments of a dielectric material such as, a thin silicon oxidelayer, prepared by chemical oxidation, for example in HNO₃, RCA, O₃, orHCl.

The wide band gap layer 10 can be tuned for enhanced or reduceddiffusion of dopant atoms from the passivation layer stack to thelight-absorbing silicon layer 1 for example by adapting its thickness,density, chemical bonding configuration and composition. It is knownthat for thermal oxidation of silicon the thickness of the grown siliconoxide layer can be controlled precisely by tuning the oxidationtemperature, time, ambient, and other process parameters. However, forchemical oxidation of silicon, for example in HNO₃, the process isreported to be self-limiting, i.e. after a certain exposure time, forexample 60 min, the thickness does not increase further.

Experiments conducted by the inventors have shown, surprisingly, thatthe thickness of a silicon oxide layer grown by exposure to HNO₃ dependson the doping concentration in the light-absorbing silicon layer 1.Choosing an appropriate doping concentration thus permits to control theoxide thickness, and thus also to control diffusion of impurities fromthe silicon-doped layer through the wide bandgap material into thelight-absorbing silicon layer 1.

Diffusion of dopants from the silicon-doped layer to the light-absorbingsilicon layer 1 through the wide band gap layer 10 can further beinfluenced for example by altering the wide band gap layer 10 by plasma,thermal, electrical or chemical treatments.

For example, thin silicon oxide layers 10 can be nitrided by exposure toa nitrogen containing plasma or by annealing in nitrogen, N₂O or NOatmosphere, which reduces diffusion of many elements through the siliconoxide layer. Further, silicon oxide layers can be densified by thermalannealing, also reducing diffusion from the silicon-doped layer to thelight-absorbing silicon layer 1. The density and composition of thebuffer layer can also be influenced by exposure to chemical agents, forexample exposing a chemical oxide grown in HNO₃ to O₃.

A given wide band gap layer 10 can also be restructured with plasmatreatments, which can result in a less dense buffer layer 4 where thelatter is present in the passivating and conductive stack PC, and whichcan enhance diffusion of dopants from the silicon-doped layer stack tothe light-absorbing silicon layer 1.

According to the invention, the wide band gap layer 10 and the thermaltreatment(s) applied to the passivating and conductive stack PC arepreferably designed such that dopants which diffuse from thesilicon-doped layer stack towards the light-absorbing silicon layer 1 donot reach the light-absorbing silicon layer 1, i.e. do not alter thedoping concentration in the light-absorbing silicon layer 1.

Furthermore, the wide band gap layer 10 may contain impurity atomsalready prior to the thermal treatment(s). During said thermaltreatment(s) these impurity atoms can then diffuse from the wide bandgap layer 10 either to the light-absorbing silicon layer 1 or to otheradjacent layers of the passivating and conductive stack PC were they canact as dopant impurities. This can be especially useful for aligning theenergy bands in the light-absorbing silicon layer 1 and in thepassivating and conductive stack PC to enhance charge carrier transportthrough the buffer layer 4. The wide band gap layer 10 can be amorphousor crystalline.

Transport through the wide band gap layer 10 can proceed via directclassical transport, or, if the buffer layer 10 is a dielectric layer,by hopping or tunnelling for example, or a combination of allmechanisms.

Part of the invention is that highly doped region 100 enhances chargecarrier transport from the light-absorbing silicon layer 1 to thepassivating and conductive layer stack. In the context of the presentinvention, “Highly doped” shall mean a concentration of the dopingimpurity of more than 10¹⁷ cm⁻³, preferably more than 10¹⁸ cm⁻³, stillpreferably more than 10¹⁹ cm⁻³. This is especially important when theelectronic band structure of the passivating and conductive stack PC isnot identical to the band structure of the light-absorbing silicon layer1. This is due to the effect that doping influences band alignment. Highdoping on both sides of a thin wide bandgap material 10 aligns themaxima of the supply functions in the layers adjacent to the dielectriclayer and thus enhances charge carrier transport through the widebandgap layer.

The terms “Supply function” shall be understood here to the chargecarrier distribution at the interface of light-absorbing silicon layer 1with the dielectric silicon-based doped layer 6 as described in[Sentaurus™ device user guide, Synopsis, Version 1-2013, Mountain View(Calif.), USA]. For this reason, the alignment effect of the supplyfunction is especially beneficial for the case of the heat-resistantheterojunction layer stack of the current invention.

Interestingly, for shallow in-diffusion, band alignment can also becaused by fixed charges instead or in addition to a highly doped region.The inventors have surprisingly found that fixed charges can cause bandbending such that passivation and/or charge carrier transport isenhanced, especially using fixed charge for influencing the alignment ofthe charge carrier supply functions. Fixed charges can be stored in anyof the layers of the passivation layer stack or at the layersinterfaces.

The layers of the passivating and conductive stack and thelight-absorbing silicon layer 1 can react with each other during thermaltreatment they undergo during their manufacturing process. This isespecially possible if adjacent layers of said passivating andconductive stack PC are made of a different materials, such as forexample SiOx and SiCx. Reaction is meant as chemical reaction, or ingeneral as any influence of one layer on the other, caused by thepresence of the former.

Avoiding such a reaction is one of the roles of the silicon-based bufferlayer 4.

Another role of the silicon-based buffer layer 4 can be to adjustdiffusion of dopants from the silicon-based doped layer 6 deposited onit towards the wide bandgap material 10 and the light-absorbing siliconlayer 1 during the at least one of the thermal treatments. Theheterojunction solar cell structure of the present invention can beapplied to different morphologies of the light-absorbing silicon layer 1such as: polished surfaces, shiny-etched surfaces, textured surfaces(alkaline, acidic or via plasma etching) and surfaces that were texturedwith subsequent surface smoothing or roughening.

Various embodiments of the inventive solar cell structure according tothe present invention are represented in FIGS. 2 to 7 as explainedbelow.

FIG. 2A represents of solar cell made from a silicon wafer 1 with alow-cost front side electrode stack (POCl₃-diffused emitter 102,SiNx-anti reflection layer 202 and low-cost fired metallization 302) andthe inventive full area heat resistant heterojunction on the rear side.The thermal treatment needed for forming the passivating contact of thisinvention can be accommodated during the thermal annealing (e.g.compatible with POCl₃, BBr₃, BCl₃ for forming the other contactpolarity) or during the firing for front side metallization.

In both cases the rear passivating contact stack consists of a thin SiOxlayer 10, an optional Si-based buffer layer 4, a doped Si-based (e.g.SiCx) layer 6 and a capping layer 12. The doped Si-based layer 6 isdesigned to be permissive to hydrogen or other passivating species suchthat the doped Si-based layer 6 is mechanically stable (i.e. does notblister) even when using firing, and a passivating species-donor layer,deposited after the firing process, does not blister upon annealing orfurther firing. Additionally, one of the layer forming the passivatingand conductive stack might contain passivating species which arereleased by the thermal treatment (firing of thermal anneal) and driventowards the interface between the wide band gap material and the Siwafer where they passivate defects.

Preferably, as shown schematically in FIG. 2A, solar cells include asurface texture. For the sake of simplicity solar cells are oftendepicted with planar surfaces as shown in FIG. 2B, as is used in theother representative figures. In any case, every surface can be texturedor planar.

A variant of part of the solar cell of FIG. 2 is shown in FIG. 5A, 5B,where the passivating and conductive stack on the silicon wafer 1comprises a wide band gap material layer 10 and doped silicon-basedlayer 6 and transparent conductive oxide layer 301 and metallizationlayer 300 (FIG. 5A), and additionally a Si buffer layer 4 (FIG. 5B).

Alternatively, to the metallization layer 300 consisting of severalelements (fingers) as shown in the FIGS. 5A, 5B it can also be appliedas full area layer (not shown).

FIG. 3A to 3C show the use of a capping layer 12 in a solar cell asproposed by the invention and shown in FIG. 2A, 2B and further in FIG.6.

A capping layer 12 is provided to ensure in-situ hydrogenation of theinterface between the light-absorbing silicon layer 1 and an adjacentlayer, the light-absorbing silicon layer 1 itself or the wide-band gaplayer 10 or the silicon-based buffer layer 4 or the doped Si-based layer6 or to redistribute hydrogen or reconfigure hydrogen bondingconfiguration in said layers. For this means, the capping layer 12:

-   -   reduces the hydrogen effusion from the layers underneath    -   preferably acts as source of a passivating specie

The capping layer 12 can be conductive and doped, and possibly be amixed-phase material. The capping layer 12 can also act as etch-stop anddiffusion barrier (e.g. Ag-glass frit penetration during firing),facilitating device integration of the passivating contact.

Indeed, some heterojunction layer stacks, for example heterojunctionlayer stacks consisting of SiOx and Si layers, are not chemically stablein hydrofluoric acid (HF). HF is used in many process steps in standardsilicon solar cell manufacturing. A capping layer with better stabilityin HF than SiOx (e.g. a Si or SiCx or SiNx layer) can thus protect theheterojunction layer during solar cell processing.

Another frequently used process in solar cell manufacturing is diffusionof impurity dopants (for example phosphorous, arsenic, boron, aluminum)to define highly doped regions in the silicon wafer. Impurity diffusionis usually carried out at elevated temperatures, typically between 700°C. and 1000° C. Frequently, the silicon wafer is exposed to a diffusionsource (e.g. POCl₃, BBr₃, BCl₃) which acts on all surfaces of thesubstrate. The diffusion source and the diffused regions in the siliconwafer need then to be removed after the diffusion process everywherethey are not wanted, typically by etching the diffusion source and alsothose parts of the silicon wafer. A more elegant solution is to apply alayer to the substrate which reduces the penetration of impurities tothe substrate, thus also called diffusion barrier. In that case, thediffusion barrier might be needed to remove afterwards, but the siliconwafer does not need to be etched.

With the capping layer acting as diffusion barrier, it protects theheterojunction layer stack during a diffusion process. Consequently, inthe cell manufacture process, it is possible to first prepare theheterojunction layer stack and then the capping layer, and then carryout a diffusion process. This is beneficial because the thermal profileof the diffusion process also acts as thermal treatment for theheterojunction layer stack.

The capping layer can be made of silicon (Si), silicon carbide (SiCx),silicon nitride (SiNx), silicon oxide (SiOx), silicon oxynitride(SiOxNy), silicon oxycarbide (SiOxCy), silicon carbonitride (SiCxNy),silicon oxycarbonitride (SiOxCyNz), which can all be doped p-type orn-type with Boron (B), Aluminum (AI), Gallium (Ga), Indium (In),nitrogen (N), Phosphorous (P), Arsenic (As), Antimony (Sb) and can behydrogenated (e.g. N-doped SiCx:H). It can also be made of a transparentconductive oxide material such as tin oxide, also doped e.g. withFluorine (F) or antimony (Sb), gallium oxide (GaOx), Indium Tin Oxide(ITO), Indium Cerium Oxide (ICO), Indium Tungsten Oxide (IWO), IndiumZinc Oxide (IZO), Zinc Tin Oxide (ZTO), indium Tin Zinc Oxide (ITZO),Indium Gallium Oxide (IGO), Indium Gallium Zinc Oxide (IGZO), Indium TinGallium Oxide (ITGO), Zinc Oxide (ZnO), Titanium oxide (TiOx), Titaniumnitride (TiNx), aluminum nitrode (AlNx), Aluminum oxide (AlOx), aluminumzinc oxide (AIZnOx), which can all contain hydrogen, Al, B, Ga, or O.The capping layer can also be a double layer of said materials, forinstance SiCx:H/SiNx:H, SiNx:H/ZTO, ZTO/ITO etc. The capping layer 12may therefore advantageously to prevent passivating species to escapefrom the buffer layer 4, doped Si-based layer 6, driving them towardsthe wide band gap material layer 10 and the wafer forminglight-absorbing silicon layer 1 where dangling bonds responsible forhigh recombination are passivated.

As shown in FIG. 3B the capping layer 12 can further serve aspassivating species donor layer in case the other layers do not containlarge amounts of passivating species, for instance after a first thermaltreatment, such as a firing treatment for example. Thus passivatingspecies is reintroduced in the Si-based doped layer 6, the buffer layer4, the interfacial wide band gap material 10 and the silicon wafer 1.

Finally, as shown in FIG. 3C the capping layer 12 can serve aspassivating species donor layer and in addition prevent passivatingspecies contained in layers 4, 6 and 10 from effusion during the firingprocess.

FIGS. 4 and 5 show embodiments of parts of solar cells according to theinvention but which do not comprise any capping layer as describedabove.

In FIG. 1A the passivating contact stack structure comprises on thelight-absorbing silicon layer 1 a wide band gap material layer 10 anddoped silicon-based layer 6 and metallization 300 for the rear backelectrode, similar to PERC cells known from prior art otherwise. FIG. 1Bfurther shows a similar structure but a Si buffer layer 4.

FIG. 4C shows a similar structure to the cell of FIG. 1A but the wideband gap material layer 10 having through holes such that the dopedlayer 6 is in direct contact with the light-absorbing silicon layer 1.Similarly, FIG. 4D shows a solar cell structure with a wide band gapmaterial layer 10 having holes through which a Si buffer layer 4 is indirect contact with the silicon wafer 1.

Alternatively to full-area the metallization can also consist of severalelements (fingers, see FIG. 2).

FIG. 6A to 6D show alternative structures of a solar cell as shown inFIG. 4A to 4D. FIG. 6A represents a passivating contact structure onwafer 1 consisting of wide band material layer 10, Si buffer layer 4,and Si-based doped layer 6 and capping layer 12 and metallization 300.In FIG. 6B a second Si buffer layer 5 is provided.

In FIG. 6C, the capping layer 12 comprises openings (holes, trenches,cracks) through which the layer on top (here a transparent conductiveoxide layer 301) is in direct contact with the doped Si-based layer 6.In FIG. 6D, the metallization 300 is filled in the openings of thecapping layer 12.

Alternatively to full-area the metallization can also consist of severalelements (fingers, see FIGS. 2 and 5).

Finally, FIGS. 7A and 7B represent rear-side contacted solar cellstructure employing a contact stacks according the present invention fordual polarities.

In FIG. 7A a light-absorbing silicon 202 passivation and anti-reflectionlayer is applied to a top surface of a silicon wafer 1 while a wide bandgap layer 10, and buffer layer 4 are provided on a bottom, rear side ofsaid wafer 1. Additionally, first electrode stack comprising a dopedSi-based layer 6 of one polarity, first TCO layer 301, first metalelectrode 300, and second electrode stack comprising doped Si basedlayer 6′ of the other polarity, second TCO layer 301′, second metalelectrode 300′ are provided to a rear side of the buffer layer 4.

The layer sequences 6, 301, 300 and 6′, 301′, 300′ form a finger gridover the rear surface, for example an interdigitated grid.

In FIG. 7B rear-side contacted solar cell as sketched to FIG. 7a butwith the doped Si-based layers 6 not being structured to accommodate thesecond polarity pillar.

Whatever embodiment of the solar cell of the invention the manufacturingthereof relies on same principles as described herein after, i.e. theprovision of thermal treatment to passivate the contact structurethrough the wide band gap material layer 10.

Several beneficial effects can be exploited for any of the wide band gaplayer 10, the silicon-based buffer layer 4, the silicon-based dopedlayer 6, and the capping layer 12 contained in the passivating andconductive stack PC of the inventive solar cell structure,

These layers 4, 6, 10, 12 can each be a graded layer, i.e. a layer whosephysical properties change along the growth direction. For example, thiscan be the crystallinity or the crystal size, the composition, thedoping, the optical properties, or the electrical properties of somecomponents or the layer as a whole. The gradient property can also becreated, or be further enhanced, during one of the thermal treatmentsundergone during manufacturing process of the inventive solar cell.

More specifically, if one or several of the layers 4, 6, 10, 12 isamorphous in the as-deposited state, layers 4, 6, 10, 12 can partiallycrystallize during the thermal treatment, starting from one side oflayers 4, 6, 10, 12, leaving the other side of layers 4, 6, 10, 12 stillmainly amorphous. The side were crystallization starts can be the sideof layers 4, 6, 10, 12 oriented towards the wafer. This effect can beexploited to reduce junction resistance between the layer and the lightabsorbing silicon wafer (improving fill factor). The other case, thelayer crystallizing only on the other side and the side oriented towardsthe wafer remaining more amorphous, can be used to obtain a morefavourable band bending thanks to band offsets between the wafer andlayers 4, 6, 10, 12, but at the same time a crystalline and thus highlyconductive nature of layers 4, 6, 10, 12 on its other side, reducingcontact resistance there. In a preferred embodiment the structure andcomposition of the heterojunction layer stack are graded over the layersuch that the refractive index is graded, the refractive index attaininghigher values towards the side oriented towards the silicon absorberwafers and lower values towards the other side. The layers 4, 6, 10, 12may contain fluorine or a fluorine compound. During a thermal treatmentthe fluorine or fluorine compound can diffuse to the interface of thesilicon wafer where it can accumulate and passivate electronic defectstates. This is especially advantageous as this passivates the interfacebetween the silicon absorber and the passivating and conductive layerstack.

For example, if any of the layers 4, 6, 10, 12 is grown as asilicon-containing layer grown by PECVD, SiF₄ can be used as precursorgas as fluorine and silicon source. To give another example, for thegrowth of boron-doped films BF₃ can be used as precursor gas tointroduce boron and fluorine. Especially advantageous is the use of aSiOx buffer layer because F accumulates at the Si—SiO₂ interface,reducing the interface defect density. FIG. 27 reveals that thermalannealing at 850° C. results in an accumulation of fluorine at the c-Siwafer and chemical SiOx interface which is beneficial to enhance surfacepassivation. At temperature of 925° C. surface passivation offluorinated passivating and conductive stack drops again due toout-diffusion of F. FIG. 28A represents cell parameters obtained usingfluorinated dopant source (BF₃) demonstrating that F can be used toobtain high surface passivation after thermal annealing and without theneed of an additional post-hydrogenation step, thus simplifying processsteps.

The layers 4, 6, 10, 12 can also be adhesion-promoting layers, alsocalled “sticky layer”.

The layers 4, 6, 10, 12 can be amorphous or crystalline. While in priorart the conductive layer is amorphous in its as-deposited state and thethermal treatment used to promote crystallization, it can also beprepared fully crystalline without any additional thermal treatment.“Fully crystalline” refers to a layer which does not crystallize furtherduring the at least one thermal treatments, i.e. does not exhibit largergrain size.

In another advantageous variant the layers 4, 6, 10, 12 can be amorphousbefore thermal treatment, and no crystallization occurs during the saidthermal treatment. Or only a fraction of it crystallizes, for instanceonly the upper part or only the lower part of the layer, leaving therespective other part amorphous, or only the intermediate part of thelayers 4, 6, 10, 12 crystallizes, leaving the upper and lower partamorphous.

Production method of an inventive heterojunction solar cell HC:

The heterojunction solar cell HC of the invention is typically preparedby cleaning a silicon wafer 1 in several chemical solutions. Typically,these are solutions in water of one of the following components H₂SO₃,HNO₃, HF, HCl, H₂O₂, O₃, NH₄OH. After most cleaning steps and after thelast cleaning step, the wafer is dipped in HF.

Then, a very thin wide band gap material layer 10 of controlledthickness, for example a SiOx layer, is grown or deposited, for exampleby exposure to hot HNO₃, UV light, H₂O₂, H₂SO₃ or O₃, or by atomic layerdeposition (ALD). Optionally a silicon-based buffer layer 4 is grown ordeposited on the SiOx layer 10.

Subsequently, a doped Si-based layer 6 is grown or deposited on thebuffer layer 4, for example by chemical vapour deposition (CVD),plasma-enhanced chemical vapour deposition (PECVD) or low-pressurechemical vapour deposition (LPCVD) or sputtering. Optionally a cappinglayer 12 may be grown or deposited then.

Afterwards, the structure is subjected to a thermal treatment treatmentat temperatures above 600° C. When combined with a front sidehomo-junction, the thermal treatment needed for forming the passivatingcontact can be accommodated during the thermal annealing for theformation of the other contact polarity (e.g. POCl₃, BBr₃, BCl₃) orduring the firing to form the Ag-metal grid.

Said capping layer 12 may remain or be removed further to the thermaltreatment, in which case a secondary capping layer may be grown,followed by an optional second thermal treatment. Any capping layer 12is prepared on top of the doped Si-based layer 6 (e.g. SiCx layer) inorder to control passivating species out-diffusion or to serve asin-situ passivating species donor layer during one of the thermaltreatment(s) (e.g. during firing or a later thermal process).

Any capping layer may preferably be able to release passivating speciesduring the following thermal treatments. passivating species aretransmitted by the doped silicon based layer 6 to the wide band gaplayer 10 and the silicon wafer 1. Eventually the first capping layer ismade such that the passivating species passes through the entire wafer1, reaching also the wafer side that is opposed to the passivatingcontact wide band gap layer stack from which the passivating species isreleased.

The doped silicon-based layer 6 is arranged advantageously to exhibitminimized mechanical stress to avoid blistering during rapid thermalannealing processes (firing). This is achieved by two means:

a. alloying Si with C, N or O or a combination thereof, binding part ofthe hydrogen to those elements instead of to the Si atoms. The bondenergy of H is higher in SiCx than in a-Si, and yet higher in SiOx, andyet higher in SiNx. This is due to the binding energy of the C—H, O—Hand N—H bonds increasing in this order. As a consequence, thehydrogen-related bonds feature a broad range of binding energies, andhydrogen evolves over a broad temperature range instead of in a smalltemperature range. This way, the evolution of free hydrogen per amountof time in the layer is minimized, which facilitates hydrogen diffusionbefore hydrogen accumulates in blisters.

b. a hydrogen-transmissive layer structure. The network structure of thelayer is capable to store or conduct hydrogen through internalnano-voids in order to avoid blistering during rapid thermal annealingprocesses (firing) and allow hydrogen reintroduction afterwards. Stressthat builds up during the thermal treatment can thus be released in-situby internal hydrogen storage and diffusion of hydrogen to othercomponents

In addition, after the thermal treatment(s) the doped silicon-basedlayer 6 exhibits very low hydrogen content (<5%).

The doping can be graded through the layer and the content of C, N or Oof the Si-alloy layer can be graded through the doped silicon-basedlayer 6. The doped Si-based layer 6 and also the Si-based buffer layer 4can be doped for example with Al, B, Ga, P, N, As, Sb, O

The doped silicon-based layer 6 can advantageously be amorphous orcrystalline, or partially amorphous and partially crystalline.

For the case too much hydrogen leaves the layer 6, it can bereintroduced after firing from an external source thanks to the Htransmissivity of the doped silicon-based layer.

The passivating and conductive stack PC of the heterojunction solar cellstructure of the invention can optionally include at least onesilicon-based buffer layer 4 between the thin wide band gap materiallayer 10 and the doped Si-based layer 6.

The manufacture of the solar cell of the invention relies essentially onthe provision of at least one thermal treatment involving temperaturesabove 600° C. In one embodiment firing are applied to form thepassivating and conductive stack PC. Such firing step might be followedby hydrogenation step involving a thermal treatment at lower temperatureor an additional firing.

In another embodiment the heat resistant passivating and conductivestack is formed by thermal annealing resulting in the formation of adiffused doped region in the light absorbing silicon with a sheetresistance between 5 and 10⁵Ω/square. In general, thermal processing atelevated temperatures causes the diffusion of dopant atoms. A centralaspect of this invention is, however, that the thermal budget and/or thepassivating and conductive stack PC, comprising a wide band gap layer 10and silicon-based doped layer 6 at least, and preferably a buffer layer4, are designed such that dopant diffusion to the light absorbingsilicon layer 1 is avoided, possibly completely omitted, creating anextremely abrupt, shallow, buried junction with the intensity of thesignal of the doping specie, expressed in [counts/s] decaying of morethan one order of magnitude within a few nm in the buffer layer 4 or thewide band gap material layer 10 or in the light absorbing silicon wafer1. This avoids possible degradation of the thin wide band gap materiallayer 10 due to dopant (for example Boron) penetration and increase ofAuger recombination or

Moreover, the different layers can crystallize partially duringannealing and the different chemical phases can be split further orintermix. More in detail, in a layer consisting of SiCx, annealing canlead to re-structuring of the layer or incorporation of C to the atomicnetwork, i.e. increasing the number of Si—C bonds. Annealing can alsolead to partial separation of SiCx to Si and SiC. Annealing can alsore-configure hydrogen bonding in the layer, for instance annealing cancause rupture of Si—H bonds but formation of C—H bonds, i.e. influencepreferential bonding of hydrogen. Annealing can also change theproperties of the buffer layer, for instance the thin silicon oxidelayer.

Annealing can also be a local process, for instance laser or microwaveannealing, i.e. exposing a local portion of the layer stack to laser ormicrowave radiation. Part of the invention is to use local annealing toform the heterogeneous layer, or to promote locally the diffusion ofdopants from the layer stack to the wafer.

In variants of the solar cell of the invention the latter comprises acapping layer, which is prepared on top of the doped Si-based layer 6(e.g. SiCx layer) in order to control hydrogen or other passivatingspecies out-diffusion or to serve as in-situ hydrogen or otherpassivating species donor layer during one of the thermal treatment(s)(during firing or thermal anneal or a later thermal process), as shownin FIG. 51. In preferred embodiments of the inventive solar cell of theinvention at least one layer of the heterojunction stack containshydrogen or other passivating species, and acts as hydrogen or otherpassivating species donor layer to passivate the interface between thewafer 1 and the wide band gap material layer 10 or the interface betweenthe wide band gap material layer 10 and the buffer layer 4 or the layerstack, or the wide band gap material layer or the buffer layer.

For the sake of completeness of the present description, table Isummarizes various preferred material for the essential layers of theheterojunction stack of the solar cell of the invention and theirpreferred thickness ranges.

TABLE I Thickness Denomination (nm) Material Wide band gap 0.5-20  SiOx,SiCx, SiNx, SiOxNy, SiOxCy, material layer (1-2)  SiCxNy, SiOxCyNz (10)AlOx, HfOx, AlHfOx, AlNx, TiNx, ZrOx, Y2Ox, AlSiOx, HfSiOx, AlHfSiOx,Buffer layer  1-100 Si; SiCx, SiNx, SiOx, SiCxNy, SiCxOy, (4) (2-10)SiNxOy, SiCxNyOz or a combination thereof Doped layer  1-300 Si, SiC,SiOx, SiCx, SiNx, SiOxNy, (6) (2-30) SiOxCy, SiCxNy, SiOxCyNz which canall be doped B, Al, Ga, In, N, P, As, Sb, and can contain hydrogen orfluorine (e.g. N-doped SiCx:H) Capping layer   2-10000 Si, SiCx, SiNx,SiOx, SiOxNy, SiOxCy, (12) (50-500) SiCxNy, SiOxCyNz, which can all bedoped B, Al, Ga, In, N, P, As, Sb, and can contain hydrogen or orfluorine (e.g. N-doped SiCx:H); SnOx,, also doped with F or Sb, GaOx,ITO, ICO, IWO, IZO, ZTO, ITZO, IGO, IGZO, ITGO, ZnO, TiOx, TiNx, AlNx,AlOx, AlZnOx, which can all contain hydrogen, Al, B, Ga, or O; metalssuch as Ti, V, Cr, Mn, Fe, Co, Cu, Zn, Y, Zr, Nb, Mo, Tc, Ru, Rh, Pd,Ag, Cd, In, Sn,, Lu, Hf, Ta, W,, Re, Os, Ir, Pt, Au, Hg, Tl, Pb, Bi, Poand rare earth metals, all of which may contain hydrogen; Metallichydrides

EXAMPLES Example 1: Boron-Doped Silicon Carbide as Passivating RearContacts Silicon Heterojunction Contact and Corresponding Solar Cell

Example 1 provides optimization of a rear hole contact for p-type solarcells which relies on full-area processes and provides full-areapassivation and conductivity. The passivating hole contact is based on achemically grown thin silicon oxide SiOx and a passivation stack ofintrinsic amorphous silicon together with in situ boron doped siliconcarbide on top, annealed at 775-900° C. The thickness of the a-Si(i)buffer layer was optimized as shown in FIG. 41. FIG. 40 shows the TEManalysis of the stack after thermal annealing. The Fourier transformanalysis in figure, clearly shows that the top SiCx(p) has lesscrystallinity compared to the SiCx(p). This is related to the fact thatadding C to the a-Si network retards its crystallization temperature.

For the optimized passivation stack, after thermal annealing, forminggas annealing increases the effective lifetime from 160 μs to 250 μs,while hydrogenation from a silicon nitride overlayer results in anincrease over 1.7 ms. A systematic contact optimization is obtained bytuning the doping concentration, annealing temperature, annealing dwelltime and monitoring the implied open circuit voltage (iVoc) and contactresistivity (ρc) in parallel. It is observed that for highly dopedlayers the optimum annealing temperature for high quality surfacepassivation is 800° C. while for lowly doped layers the optimumannealing condition shifts to 850° C. Excellent surface passivation andefficient current transport is evidenced by an iVoc value of 718 mV anda ρc of 17 mΩcm² on p-wafers.

Proof of concept p-type hybrid solar cells with a standardheterojunction front contact prove the excellent efficiency potential ofthe passivating rear contact by reaching a Voc of up to 709 mV and a FFof up to 80.9%. The best hybrid cell achieves a conversion efficiency of21.86%, enabling Voc of 708 mV, FF of 79.9% and Jsc of 38.7 mA/cm² afterannealing at 825° C.

Experimental Framework

Passivating hole contacts were first investigated using symmetricalstructures based on 250 μm thick chemically polished 4-inch p-type floatzone <100> c Si wafers with a resistivity of 2 Ωcm.

After standard wet chemical cleaning, an ˜1 nm SiOx wide band gap layerwas formed by wet chemical oxidation using HNO3 solution at 80° C.,referred to as “chemical SiOx” hereinafter. Subsequently, a layer stackconsisting of a 10 nm thick intrinsic silicon [Si(i)] interlayer and a30 nm thick in situ boron doped silicon rich silicon carbide [SiCx(p)]layer was deposited on both sides by plasma enhanced chemical vapourdeposition (PECVD), both layers being amorphous in their as depositedstate, i.e. before any thermal treatment.

Following the PECVD step, the samples were annealed in inert gasatmosphere at temperatures between 775° C. and 900° C. with differentannealing dwell times using heating and cooling ramps of 10 and 2°C./min, respectively.

This was followed by a hydrogenation process to passivate electronicdefects at the Si wafer/chemical SiOx interface, either by annealing at500° C. for 30 minutes in forming gas (4% H2 in N2) (FGA) or by applyinga silicon nitride (SiNx) overlayer as a hydrogen source. For the latter,samples were annealed on a hot plate at 450° C. for 30 minutes torelease hydrogen from the SiNx, which is then followed by etch backprocess in HF to remove this overlayer. For the SiNx deposition, a VHFPECVD tool at the excitation frequency of 81 MHz was used and thesubstrate temperature was set to 250° C. during the deposition.

The effective minority-carrier lifetime was investigated by photoconductance decay (PCD) technique as a function of the excess minoritycarrier density giving access to the implied open circuit voltage (iVoc)values. The contact resistivity (ρc) was measured using the transferlength method (TLM) after sputtering ITO/Ag contact pads through ashadow mask on p-type wafers. The samples were dipped in 5% diluted HFfor 1 minute right before the sputtering process to remove the surfaceoxide. The spatial homogeneity of the passivation was analyzed usingphotoluminescence imaging (PLI). For thickness determination, thedeposited layers were characterized using spectroscopic ellipsometry(SE, HORIBA Jobin Yvon, UVISEL). To measure the hydrogen effusion as afunction of temperature thermal desorption spectroscopy (TDS) wasperformed in an EMD□WA1000S/W system at the National Institute ofAdvanced Industrial Science and Technology (AIST), Japan. For thismeasurement, only one-side deposited samples were prepared. Before themeasurements samples were kept one hour at room temperature under lowpressure to degas as much as possible the water absorbed on the c-Siwafer surface and the chamber walls to reach a base pressure of5.25*10⁻¹⁰ Torr.

For the measurement, samples were heated with a halogen lamp from roomtemperature to 690° C.-substrate temperature-, at a rate of 20° C./min.Effusing species were identified by quadrupole mass spectrometry. Theelectrochemical capacitance-voltage (ECV) profiling technique wasemployed to measure the diffusion profiles of boron within the Si-waferfrom the SiCx(p) layers using 0.1 molar ammonium fluoride (NH4F)solution as etchant.

For proof of concept hybrid solar cells, the passivating hole contact onthe planar rear side of a single-side textured wafer was first prepared.Then a silicon heterojunction electron selective front contact made ofintrinsic amorphous silicon (a-Si:H) and phosphorus-doped a-Si:H(n)layers was built. Finally, 70 nm and 200 nm indium tin oxide (ITO) weresputtered onto the front and rear sides of the cells, respectively. Theactive cell area was defined by depositing the ITO layers through a 2.2cm×2.2 cm shadow mask. A silver reflector/contact was sputtered onto therear side and silver paste was screen printed to realize the frontmetallization grid, followed by curing for 30 minutes at 210° C. in abelt furnace. Detailed information about the fabrication process for thehybrid cells can be found in G. Nogay, J. Stuckelberger, P. Wyss, Q.Jeangros, C. Allebé, X. Niquille, F. Debrot, M. Despeisse, F.-J. Haug,P. Löper, Silicon-Rich Silicon Carbide Hole-Selective Rear Contacts forCrystalline-Silicon-Based Solar Cells, ACS Applied Materials &Interfaces, 8 (2016) 35660-35667. Current voltage (I-V) characteristicsof the cells were measured at 25° C. with a source meter (Keithley,2601A), using an AAA solar simulator (Wacom) calibrated to 100 mW·cm⁻²with a c-Si reference cell. Additionally, the suns-Voc method was usedto obtain the series-resistance-free pseudo I-V curves.

Results

FIGS. 8A and 8B show the evolution of the effective lifetime values ofsymmetric sister samples after annealing for boron diffusion, afterforming gas annealing (FGA) and after application of the different SiNxlayer as a hydrogen donor layer, considering (A) the effect of theNH3/SiH4 flow ratio during the SiNx deposition, and (B) the effect ofthe SiNx thickness when layers are deposited with the NH3/SiH4 flowratio of 1.9.

Most passivating contacts based on thin silicon oxide layers prepared bychemical passivation require reintroduction of hydrogen after thehigh-temperature annealing step in order to decrease the interface trapdensity and attain good surface passivation. FGA is one of the mostrecognized methods for re hydrogenation of such interfaces andimpressively high iVoc values have been reported for n-type poly-Sicontacts on n-type Si wafers using this method.

According to the inventive structure annealing in forming gas indeedimproves the passivation properties, as evidenced by an increase of theeffective lifetime from 160 μs to 250 μs, as seen in FIG. 8A. To explorealternative hydrogenation methods the same samples were used anddeposited a SiNx layer as hydrogen source. For the latter, thedeposition parameters of the SiNx layer were tuned to improve thehydrogenation quality i.e. the passivation quality. FIGS. 8A and 8B showthe evolution of the effective lifetime values along the contactformation process for an inventive contact stack prepared with a TMBflow of 1.5 sccm, at the injection level of 10¹⁵ cm⁻³. After diffusionannealing we observe rather low effective lifetimes around 160 μs. Afterannealing in forming gas at 500° C. for 30 minutes an improvement in thelifetime is observed, which is, however, is a moderate increase incontrast with other reported values for forming gas annealed borondiffusions passivated by chemical SiO₂. Following the FGA, a 60 nm thickPECVD SiNx layer, produced with a NH₃/SiH₄ flow ratio of 0.5, wasapplied to both side of the symmetric test structures to investigate theeffect of the atomic hydrogenation.

FIG. 8A shows an impressive improvement of the effective lifetime afterhot plate annealing at 450° C. for 30 minutes. Note that these valueswere measured after etching the SiNx with 5% diluted HF solution,ensuring that the observed effect does not originate from field effectpassivation caused by fixed charges in SiNx. Consequently, SiNx acts asa hydrogen reservoir and hot plate annealing induces the hydrogendiffusion through the p contact stack to the chemical SiOx/c Si waferinterface.

To investigate the effect of the NH₃/SiH₄ flow ratio during the SiNxdeposition on hydrogenation at 450° C., the ratio was changed from 0.5(refractive index of 2.5 at 635 nm) to 3 (refractive index of 1.8 at 635nm). It was found that the optimal ratio of 1.9 resulted in the highesteffective minority carrier lifetime for the investigated annealingconditions. It should be noted that the optimum annealing temperaturefor the layer produced with the ratio of 3 might be higher than 450° C.since with increased NH₃ flow, density of the strong N—H bond increases.The SiNx hydrogen donor layer was further optimized by investigating theeffect of the thickness, as displayed in FIG. 8B. With increased layerthickness, first the lifetime value increases which can be attributed tothe augmented hydrogen reservoir. This trend is then followed by asaturation at around 72 nm and a drop when the layer thickness reachesto the 105 nm, however; the reason for this drop is not yet clear.

FIGS. 9A and 9B show the TDS spectra of the desorbed molecular hydrogen,H₂ (M/z=2) and water, H₂O (m/z=18) from (i) an as deposited a SiCx(p)layer, (ii) a SiNx layer deposited directly on bare wafer, and (iii) aSiNx layer deposited on a SiCx(p) layer that was previously annealed at800° C.

FIG. 9A, 9B shows that the sample with as-deposited a SiCx(p) contactlayer effuses hydrogen already at temperatures above 150° C. Theeffusing species are dominantly H₂ molecules, showing a high effusionspectra in the temperature range of 200-500° C. (FIG. 9A). The H₂(M/z=2) effusion is, in fact, almost two orders of magnitude higher thanthat of H₂O (FIG. 9B). The origin of H₂ effusion from the as-deposited aSiCx(p) is linked to the high hydrogen content in the layer [Beyer,Journal of Non-Crystalline Solids, Hydrogen evolution from a-Si:C:H anda-Si:Ge:H alloys, 1983]. The H₂ desorption spectra of the SiNx on barewafer and deposited on SiCx(p) shown in FIG. 9B reveals thatirrespective of the substrate, the observed H₂ effusion shows the samecharacteristic up to the 300° C. which is close the SiNx depositiontemperature. The high H₂ effusion, but slower desorption rate (comparedto the SiCx(p)), confirm the role of SiNx as a hydrogen reservoir. Above300° C., the layer on the SiCx(p) effuses more H₂ than the layer on thebare wafer. This can be explained by the additional H₂ in the SiCx dueto its hydrogenation during SiNx growth, or with a lower diffusivity ofH₂ in the SiCx layer than the silicon wafer, the SiCx layer thusblocking effusion of H₂ from the SiNx layer towards the silicon wafer.

The observed flat plateau after 475° C. for SiNx deposited on annealedSiCx(p) contact and after 500° C. for SiNx on bare wafer can beinterpreted as an artefact originated by the H₂ saturation in thechamber which cannot be pumped out after certain threshold.

The effect of boron doping concentration in the SiCx(p) layer on thesurface passivation and contact resistivity was investigated as afunction of annealing conditions. FIGS. 10A to 10C show the iVoc valuesfor samples produced by varying the trimethylboron (TMB) flow during thedoped layer deposition from 0.2 sccm to 1.9 sccm and annealed attemperatures range of 775° C. to 900° C. with various annealing dwelltimes. Annealing dwell times are reported as the total time thetemperature is within +/−5° C. of the dwell temperature.

For TMB flows of 1 sccm and 1.5 sccm, optimum passivation is attainedannealing at 800° C. with a dwell time of 5 minutes, leading to an iVocof 718 mV. Whereas, for lower TMB flows such as 0.2 sccm and 0.5 sccm,optimum passivation is reached annealing at 850° C. for 0 minute dwelltime, which leads to an iVoc value around 716 mV. This shift in optimumcondition can be interpreted by the fact that for decreased dopingconcentration, a higher annealing temperature is required to obtainboron diffusion through the chemical SiOx to strengthen the field effectpassivation. For the samples produced with high TMB flows such as 1.5sccm and 1.9 sccm, the dwell time does not make a significant differencein iVoc when the samples are annealed at 800° C. and 825° C., whereasfor the annealing temperature of 850° C., 15 minutes annealing dwelltime degrades the passivating quality drastically. For this dwell time,the longest used in this series, the best performance is obtained forsamples produced with lower TMB flows-except the sample with 0.2 sccmand annealed at 800° C. (see FIG. 10C). It is also perceived that with900° C. annealing, regardless of the annealing dwell time, the iVocdrops drastically, possibly due to the strong boron in diffusion whichdeteriorates the chemical SiOx/c Si wafer interface and/or due totemperature induced local disruption of the chemical SiOx. We note thatthe annealing conditions marked with an asterisk were used forprocessing the hybrid cells presented below.

In FIGS. 10A to 10F the same annealing temperatures are indicated withthe same symbols while the symbol fillings are changing with dwelltimes.

Specific contact resistivities (ρc) were characterized using the TLMmethod. FIGS. 10D to 10F show the ρc values for the samples produced byvarying the trimethylboron (TMB) flow during the doped layer depositionfrom 0.2 sccm to 1.9 sccm (sccm standing for Standard Cubic Centimetersper Minute, as defined at Tn=0° C. or 273K, Pn=1.01 bar) and annealed attemperatures of 800° C., 850° C. and 900° C. with annealing dwell timesof 0, 5 and 15 minutes. For all investigated dwell times, ρc decreaseswith increased annealing temperatures as expected. This trend can beexplained with stronger doping of the wafer next to chemical SiOx by indiffusion of boron from the SiCx(p) layer. For sufficiently highannealing temperature, it was also reported that the chemical SiOx mightbecome disrupted, leading to local thickness variation or pinholeformation which may also provoke the low ρc. Even though there is somediscrepancy for some of the samples prepared with TMB flow of 1.5 sccmand annealed at 800° C. for 0 minute (FIG. 10D) and sample prepared withTMB flow of 2 sccm and annealed at 800° C. for 15 minutes (FIG. 10F), ingeneral those figures show that with increasing boron doping density,the ρc decreases irrespective of the annealing dwell times. Thismentioned discrepancy observed in some of the samples, might originatefrom processing inaccuracies during the sputtering through the shadowmask, sometimes resulting in inhomogeneous contact pad widths, andtherefore improperly defined pad distances. The lowest ρc reached was6.84 mΩcm² with the sample produced with TMB flow of 1 sccm and annealedat 9000° C. for 5 minutes (see FIG. 10E), however, as seen in FIG. 10B,this sample shows very poor passivation quality with iVoc of 650 mV. Thehighest ρc measured is 97.26 mΩcm² with the sample produced with TMBflow of 0.2 sccm and annealed at 800° C. for 0 minute (see FIG. 10D).According to FIG. 10A, this sample has a iVoc of 712 mV. The optimalcondition for both good passivation quality and low ρc is defined withthe sample produced with 1.5 sccm TMB flow and annealed at 800° C. for 5minutes leading to iVoc of 718 mV and ρc of 17 mΩcm². Hence, the TLMmeasurements show that the presented layer stack provides not only anexcellent surface passivation but also a promising contact resistivityfor a sufficient transport of the majority carriers for full-areacontacts.

The presence of dopant diffusion through the thin chemical SiOx withinthe wafer is essential to selectively allow the transport of one type ofcarriers. The amount of dopants together with diffusion depth definesthe carrier selectivity of the contact. FIG. 11 shows the dopingprofiles within the p-type c-Si wafer measured by ECV for differentannealing dwell temperatures and dwell times for the contact that wasproduced with the TMB flow of 1.5 sccm on top of chemical SiOx and 10 nmthick intrinsic Si interlayer. Within a surface region of about 30 nmfrom the chemical SiOx/c-Si wafer interface (not shown here), theyreveal a similar behavior for all investigated temperatures with somefluctuations.

These fluctuations in the highly doped contact layers might indicate thefact that the annealing step transformed the layers into polycrystalline material where high impurity concentrations atgrain-boundaries and lower concentrations within the grains are probedsimultaneously. Beyond that region, they show the expected behavior ofdeeper in diffusion at higher temperatures, reaching a depth of 70 nmafter the annealing at 900° C. As it is expected the sample annealed at800° C. for 0 minute dwell time shows the lower carrier concentrationwithin the wafer and very shallow diffusion profile.

Additionally, the effect of annealing dwell time for 850° C. annealingis revealed in FIG. 11, and it is observed that with increased dwelltime from 0 to 15 minutes, diffusion depth increases approximately 10nm. Overall, the diffusion profiles in the bulk presented here arecomparable with the literature values obtained by applying in-situ borondoped deposited layer as a dopant source, however, distinctively theyshow a drop with an extremely steep tail to the bulk concentration ofca. 10¹⁶ cm⁻³.

Finally, the hole-selective contact was integrated in full devices toinvestigate their performance at device level. To this end, thehole-selective junction was prepared on the rear side and finished thesolar cells with a SHJ front side as described above. The usage of aheterojunction front side is motivated by assuming that the Voc of thecells will be limited by the rear side.

The obtained output characteristics for the hybrid cells exhibiting theboron doped layer produced with the TMB flow of either 1.5 sccm or 1.9sccm and annealed at dwell temperatures of 8000° C., 825° C., and 850°C. are presented in FIGS. 12A-12D. The applied annealing conditions arealso indicated in FIGS. 12A-12D to compare the iVoc values of thesymmetric test samples. FIGS. 12A-12D reveals that the samples producedwith 1.5 sccm of TMB flow shows approximately 10 mV higher iVoc valuescompared with the ones produced with 1.9 sccm for all the annealingconditions performed to realize the hybrid cells.

FIG. 12A shows that Voc values are mostly following the presented iVoctrend. For the annealing temperature of 800° C. and 850° C., the Voc ofthe cells with 1.5 sccm are higher, while for the 825° C. both dopingconditions exhibit almost the same Voc values. Overall, for the cellsannealed at 800° C. and 825° C., the difference in between iVoc and Vocis changing from 2 mV to 9 mV, for the cells annealed at 850° C. thisdifference is higher than 20 mV. This variation in between iVoc and Vocarises mostly from the ITO sputtering damage which was also reported byother groups. The best Voc reached is 709 mV with the cell produced byusing 1.5 sccm TMB flow and annealing at 800° C.

According to FIG. 12B, the obtained Jsc values are higher for thesamples produced with 1.5 sccm TMB flow. Since all the hybrid cellspresented here have similar front sides, we presume that the differenceis associated with the free carrier absorption which is greater for thesamples with higher doping concentration. For the sample with 1.5 sccmTMB flow and annealed at 825° C., the Jsc is much higher compared to therest. This might be due to the fact that even though we used the samesputtering parameters for the front ITO depositions, we observe somefluctuations in the layer properties from one deposition to another. Itcan be possible that front ITO is more transparent for this specificcell.

FIG. 12C represents the FF values as a function of annealing temperaturefor two different TMB flows. It is observed that the cells produced withTMB flow of 1.9 sccm show higher FF for all annealing conditionscompared with the 1.5 sccm. For both TMB flows, the FF increases withannealing temperature, in agreement with the trends observed from TLMcharacterization. The best FF reached is 80.9% with the cell produced byusing 1.9 sccm of TMB flow and annealing at 850° C. Overall, the cellsproduced with 1.5 sccm and 1.9 sccm show similar efficiency, especiallyfor the annealing temperatures of 800° C. and 850° C., since the lowerVoc and Jsc values for the hybrid cells produced with 1.9 sccm TMB flowis compensated with the FF gain. The best hybrid cell from this seriesreaches a conversion efficiency of 21.86%, enabling Voc of 707 mV, FF of79.9% and Jsc of 38.7 mA/cm2.

In conclusion, the present invention provides a detailed optimization ofthe thermally stable hole selective rear contact for p-type c-Si solarcells prepared by PECVD with in situ doping followed by thermalannealing. The hydrogenation process is of utmost importance to obtainan optimized chemical SiOx/c-Si wafer interface and a high-quality highlow junction after the thermal process necessary for dopant diffusionand recrystallization of the PECVD layer. In addition hydrogenationprocess using SiNx overlayer is more efficient than the FGA and it leadsa drastic improvement in lifetime for the presented SiCx(p) basedpassivating contact. With symmetric lifetime test sample, an iVoc valuesof 718 mV and ρc of 17 mΩcm² is reached after annealing at 800° C. Theproof of concept hybrid solar cell, which features the passivating rearcontact developed here and a heterojunction front side on a p-typewafer, yields an efficiency of 21.86%, enabling Voc of 708 mV, FF of79.9% and Jsc of 38.7 mA/cm².

Replacing TMB with BF3 as gas precursor during the deposition of thesilicon-based doped layer was found to be a valid solution to avoidhydrogen reintroduction, thus simplifying process fabrication. As shownin FIGS. 25 and 25 B flow ratio between CH4/(CH4+BF3/SiH4) of 0.1 wasfound as optimal value between i-VOC and contact resistivity (ρc). FIGS.26 and 28A demonstrate that high surface passivation can be obtainedafter thermal annealing without the need of an additionalpost-hydrogenation step. FIG. 27 reveals that thermal annealing at 850°C. of BF3 doped SiC(p) results in an accumulation of fluorine at thec-Si wafer and chemical SiOx interface which is beneficial to enhancesurface passivation. At temperature of 925° C. surface passivation offluorinated passivating and conductive stack drops again due toout-diffusion of F.

Example 2: Implementation of a Wide-Band Gap Silicon Carbide Layer forFront Side Carrier Selective Contacts

In this second example phosphorous doped silicon carbide (SiCx(n))layers grown on a chemically oxidized textured silicon wafer areconsidered as window layers for front side high temperature carrierselective contacts. Implied open circuit voltage higher than 735 mV wereachieved for carbon-rich SiCx(n) on textured surfaces.

FIGS. 13A, 13B and 13C show photoconductance (PCD) measurementsperformed with Sinton instrument of recombination current density (J0)(left axis), estimated with Kane Swanson method at injection levels of10¹⁶ cm−3 and i-VOC (right axis) as function of the normalized to themaximum gas flow ratio defined as the r=CH4/(CH4+PH3+SiH2+H2) of SiCx(n)layers deposited via PECVD on p-type float zone (FZ) double sidetextured, silicon wafers (sketched in FIG. 13D) with thickness around250 μm and resistivity between 1-5 Ω cm. The SiCx(n) layer thickness ˜13nm was measured with ellipsometer on polished samples (deposition timescaled on 1.7). A SiOx ˜1.4 nm thick was chemically grown prior thePECVD deposition. J0 and iVoc (as function of r after annealing at (A)800° C. for 5 min and (B) 850° C. for 5 min and (C) 900° C. for 5 min.The empty symbols indicate J0 and iVoc after hydrogenation performed byusing a SiNx as H-donor layer.

The figures clearly shows that, J0 increases after thermal annealingwith increasing of r. While after hydrogenation J0 decreases withincreasing of r, for both thermal annealing at 800 and 850° C. Thepost-hydrogenation has only a small impact on J0 for the sample with lowr but drastically reduces it for the samples with high r. This suggeststhat recombination rate of the sample with low r is mainly determined byrecombination in the in-diffused region. For samples with high r, whichare expected to be more defective, hydrogenation enhanced chemicalpassivation at Si/SiOx interface and passivates layer defects. TheSiCx(n) with r=0.8 annealed at 800° C. for 5 min lead to the lowest J0,of ˜5 fA/cm2 and highest iVoc of 735 mV. To further enhance layercrystallization higher annealing temperature would preferable. For thesamples annealed at 900° C. for 5 min regardless r and hydrogenationhigh J0 were observed. This would indicate that the recombination rateis dominated by high doping concentration from in-diffusion of P dopantatoms during annealing and/or recombination at Si/SiOx interface due tooxide breakage (TEM and ECV measurements will be performed).

The impact of the PH3 flow on the passivation quality for a givenC-content at annealing temperatures of 800° C. and 900° C. for 5 min areshown in FIGS. 14A-14B, respectively for similar SiCx(n) layersdeposited via PECVD on p-type float zone (FZ) double side texturedsilicon wafers (sketched in FIG. 13C).

All SiCx(n) layers were deposited on top of chemically grown SiOx ˜1.4nm thick. The SiCx(n) layer thickness ˜15 nm was measured withellipsometer on polished samples (the deposition time was scaled on1.7). Teff (and i-VOC as function of PH3 flow after annealing at (a)800° C. for 5 min and (b) 900° C. for 5 min. Empty symbols representsTeff and i-VOC the after hydrogenation performed via forming gasannealing (FGA) and using SiNx as H-donor layer. For both annealingconditions τeff increases towards higher PH3 flow. This indicates thatpassivation due to the in-diffused region improves with increasing ofthe doping level. Once again, the impact of the hydrogenation on Teff isvery strong for the samples annealed at 800° C. but less visible for theones annealed at 900° C. even at low PH3 flow. Together with theobservation that the passivation level of the samples annealed at 900°C. is much lower than of the 800° C. annealed samples, this hints atinterface recombination due to SiOx breakage at 900° C. rather thanrecombination in the in-diffused region (TEM measurements are underinvestigation).

The optical and the electrical properties of the SiCx(n) contact wereevaluated using it as front emitter in a p-type c-Si solar cells. Asback surface field (BSF), a Si-rich boron doped SiCx(p) doped with BF3optimized on chemically polished surfaces was used. The main solar cellfabrication steps involved: wafer cleaning, chemical oxidation, PECVD ofSiCx(n) (front) and SiCx(n) (rear), co-annealing at 800° C. for 5 min,hydrogenation and metallization. The τeff as function of the minoritycarrier density is reported in FIG. 16 for the SiCx(n) deposited ondouble side textured wafer (FIG. 15A), SiCx(p) deposited on double sidechemically polished wafer (FIG. 15B) and SiCx(n) on the textured frontand SiCx(p) on the polished rear side of the solar cell precursor (seeFIG. 15C). The figure also reports the iVoc and implied fill factor(i-FF) for all three cases. The solar cell precursor showed an iVoc of710 mV and i-FF of 85%.

A solar cell demonstrator was prepared as presented in FIG. 17. Thedemonstrator structure is a typical p-type c-Si solar cell employing theC-rich SiCx(n) (r=0.8 and r=1) on the front side and Si-rich SiCx(p) BF3doped on the rear side deposited. Both layers were deposited via PECVDon top of chemically grown SiOx layer. Co-annealing at 800° C., 850° C.and 900° C. for 5 min followed by hydrogenation via forming gasannealing and with SiNx as hydrogen donor layer was used. Finally,indium tin oxide (ITO) was sputtered on both sides of the wafer followedAg sputtering on the rear and screen-printing of Al on the front side.

Solar cells employing co-annealing process based on slow annealingtreatments for processing of both passivating contacts on the front andthe rear side were prepared, based on SiCx(n) and SiCx(p) respectively.SiCx layers were deposited via plasma enhanced chemical vapourdeposition (PECVD) on chemically oxidized Si surfaces of a siliconwafer, and are in-situ doped with PH3 (SiCx(n) with r=1) and BF3(SiCx(p)) during the deposition process.

Subsequent to deposition, annealing in a tube furnace is performed inorder to diffuse dopants and passivating species into the wafer andcrystallize the deposited layers. After thermal annealing anhydrogenation step comprising the deposition of a SiNx as H-donor layerfollowed by an annealing at 450° C. for 30 min is performed.

Fill factor (FF) up to 84% and open circuit voltage (VOC) up 728 mV weredemonstrated on planar surfaces as reported in FIGS. 28A and 28B. Thisindicated that high r value (C-rich layers) do not limit charge carriertransport. For front side textured solar cells the similar process asfor the planar cell fabrication was used. For the SiC(n) we tested twodifferent r values of 0.8 and 1. For the r=0.8, VOC up to 710 mV and FFup to 80% with a final efficiency of 21.5% were reached, see FIG. 17. Itwas however observed that parasitic absorption in the SiCx(n) frontcontact limits the potential of such solar cells, as C-addition to theSilicon network increases the temperature at which layer crystallizationoccurs.

As shown in FIGS. 18A and 18B, layers with lower flow ratio (presumablylow C-content) exhibited an external quantum efficiency which increasein the wavelength range between 300 and 500 nm with increasing of thethermal annealing temperature. Such enhanced spectral response isassociated to an increase for the layer crystallinity with increasing ofthe temperature of the thermal treatment. On the contrary the solarcells with high normalized flow ratio exhibited a EQE response in thewavelength range between 300 and 500 nm independent by the annealingtemperature. This means that increasing of the C-content of the dopedSiCx(n) layer leads to an increase of the temperature at whichcrystallization occurs leaving the layer amorphous and thus with anindirect bandgap. Such limitation can be tackled by implementation of amicro-crystalline growth regime of the SiCx(n) layer on the frontcontact. The micro-crystalline silicon carbide (μc-SiCx(n)) layer showsreduced optical absorption, demonstrating the potential of the approach.This means that layers are crystalline already in the as-deposited stateand that the thermal treatment further increases layer crystallinity.

To further decrease parasitic absorption (especially at shortwavelengths) of the front side contact we are investigating PECVDdeposited phosphorous doped micro-crystalline μc-SiCx:P layers. Having acrystalline layer structure already in the as-deposited state wouldstrongly foster full crystallization even in case of moderate annealingtemperature of 700° C. to 900° C.

FIGS. 19A and 19B show PCD measurements of τeff at injection level of10¹⁵ cm−3 (left axis) and i-VOC (right axis) as function of theannealing temperature (dwell time constant 5 min) of PECVD μc-SiCx:P(r=0.1) and PECVD SiCx:P layer (r=0.8) deposited on chemically oxidizedc-Si p-type FZ wafers (thickness 280 μm and resitivity between 1-5 Ω·cm)double side polished, detailed in FIGS. 19C and 19D, respectively. Thelayer thicknesses measured with ellipsometer are also reported. Filled(empty) circles indicate τeff and i-VOC after annealing (hydrogenation).Hydrogenation was performed via forming gas annealing (FGA) and using aSiNx as H-donor layer. For the SiCx:P (r=0.8) the highest bestpassivation was achieved with annealing at 800° C. for 5 min, τeff>8 msand i-VOC of 740 mV. For the μc-SiCx:P (r=0.01) the highest performancewere achieved after annealing at 800° C. for 5 min, τeff>0.5 ms andi-VOC of 690 mV. The impact of the hydrogenation is very strong forSiCx:P at both 800 and 900° C. For the μc-SiCx:P the hydrogenreintroduction leads to a small improvement of the surface passivation.This is most luckily related to the stronger P-diffusion for μc-SiCx:Pcompared with SiCx:P limiting the overall recombination rate.

The refractive index spectra (n,k) of the μc-SiCx:P layer afterannealing are reported in FIG. 20A-20B, which respectively represent thereal and imaginary parts of the complex refractive index measured forSiCx:P with r=0.8 annealed at 800° C. (squares), 900° C. (circles) andμc-SiCx:P annealed at 800° C. (triangles) and 900° C. (diamonds). Thesegraphs clearly shows that higher annealing temperature improve layercrystallinity decreasing parasitic absorption.

Example 3: Locally Conductive Transport Channel Formation in HighTemperature Stable Hole Selective Silicon Rich Silicon CarbidePassivating Contact Stack

This third example provides a detailed optimization and characterizationof a high temperature stable hole selective passivating contact based ona chemically grown thin silicon oxide (chem-SiOx) and boron dopedsilicon carbide (SiCx(p)) layer prepared by plasma enhanced chemicalvapour deposition (PECVD). The developed contact can be implemented atthe rear side of industrial p type solar cells to update existingmanufacturing lines. Excellent surface passivation and efficient currentextraction were observed.

A proof of concept p-type hybrid solar cell featuring this passivatingrear contact and a standard heterojunction front side was prepared (seeFIG. 21). Conductive atomic force microscopy (c-AFM) was used in orderto elucidate whether majority carrier conduction proceeds locally(possibly through pinholes or local disruption of the chem-SiOx formedin the SiOx during thermal treatment), or over the full area bytunneling. The c-AFM analysis reveals the formation of non-uniformlydistributed locally conductive areas caused by the thermal process.

Additionally, it is observed that increasing the annealing temperatureincreases the density of these conductive areas, which correlates todeteriorated passivation quality. These findings present strongexperimental evidence that the SiOx buffer layer becomes disrupted,leading to local SiOx thickness variation. C-AFM current maps taken at−100 mV with for the sample prepared with TMB flow of (a) 0.2 (b) 0.5,(c) 1 and (d) 1.5 sccm, annealed at 850° C. show that with increasingdoping concentration (FIG. 22A-22D), density of the conductive spots areincreasing, which correlates with deteriorated passivation quality anddecreased specific contact resistivities. FIG. 23A to 23D further showc-AFM current maps at −100 mV for the samples deposited with the TMBflow of 1.5 sccm (a) as deposited, (b) annealed at 800° C., (c) 850° C.and (d) 925° C. The average current is again increasing with increasedannealing temperature together with density of the conductive spotsrepresented with black color. For annealing temperature above 925° C. asevere drop of the surface passivation quality occurs (iVoc of 645 mVagainst iVoc of 715 mV for the samples annealed at 800° C.). TEM and EDXmap on FIG. 24 A of the samples annealed at 850° C. shows no disruptionof the interfacial SiOx. On the contrary, as FIG. 24 B indicates thesample annealed at 925° C. exhibits the formation of local breakage ofthe tunneling SiOx (wide bandgap layer) also known as pinhole.

Example 4: c-Si Solar Cells with Fired Silicon-Based Heterojunction asHole Collector—FLASH Heterojunction

A fired silicon heterojunction contact (FlaSH) according to theinvention and SHJ with open circuit voltage above 700 mV and conversionefficiency above 21% was prepared and tested.

This novel contact scheme allows overcoming practical integration issues(tailored annealing process or thermal instability) and physicallimitation (Auger recombination and parasitic absorption of thein-diffused region) of existing carrier selective contacts, representinga cost-effective solution for next generation of high efficiencyindustrial c-Si solar cells.

The FlaSH contact is fabricated according to sequence described in FIG.29. An ultra-thin (˜1.4 nm) SiO_(x) layer is wet chemically grown andcapped with a-SiCx(p):H deposited via plasma enhanced chemical vapordeposition (PECVD).

The FlaSH Concept:

The FlaSH contact described in the following is fabricated according tothe process sequence depicted in FIG. 29. After wafer cleaning, a thinSiOx is chemically grown followed by plasma enhanced chemical vapordeposition (PECVD) of an a-SiCx(p):H or alternatively aa-Si(i):H/a-SiC(p):H layer stack. The overall structure undergoes afiring and hydrogenation process. We used PECVD as deposition techniquefor the FlaSH contact as it offers a multitude of possibilities toprepare layer stacks with unique properties, more difficult to achievewith other techniques (e.g. with low-pressure chemical vapor deposition(LPCVD) of polysilicon followed by ion implantation or dopingdiffusion). Having tunable properties of the layer stack opens the wayfor the design of contact solutions that, for instance, can undergopre-established thermal processes, such as firing. Additionally, thisdeposition technique enables single side wafer processing and in-situdoping introducing large flexibility in terms of process integration. Arobust mechanical stability against fast thermal processes (e.g. 50°C./s) at elevated temperature (e.g. 800° C.) typical of firing definesthe first requirement for the FlaSH contact. In fact, during the firingprocess, fast hydrogen effusion process and thermal stress in the a-Si:Hlayer can result in layer blistering and delamination. Owing to thehigher bonding energy of the C—H compared to the Si—H bonds, hydrogeneffusion is retarded to higher temperatures as carbon is added to theamorphous silicon network. Consequently, SiCx layers are more resilientagainst blistering when deposited on SiOx compared to a-Si:H, hencemotivating their employment as carrier selective layers in the FlaSHstack. Additionally, SiCx films are chemically stable to wet solutionscommonly used for c-Si solar cell processing. The short duration of thethermal process leads to a partial re-crystallization of the depositeda-SiCx(B):H into a nano-crystalline (nc) nc-SiCx(p), i.e. crystallinesilicon grains in the range of few nanometers to few micrometers withinthe amorphous silicon phase, and its doping activation.

Full-area charge carrier selectivity in “conventional” intrinsic/dopeda-Si:H heterojunction contacts, is achieved via saturation of surfacedangling bonds by hydrogen and the amorphous silicon network and bandbending in the c-Si wafer induced by the highly doped a-Si:H layerresulting in a depletion of minority charge carriers from the Sisurface. The former passivation mechanism is known as chemicalpassivation and is typically expressed in terms of defect density ofstates (Dit) at c-Si/passivating layer interface. The latter is calledfield effect passivation and is set by the work function differencebetween the highly doped a-Si:H layer and the lightly doped c-Si wafer.When exposed to temperatures above 250° C., hydrogen effusion from thea-Si:H leads to a strong degradation of the chemical passivation 16.Surface passivation can be recovered (or at least partially) byemploying hydrogenation process if the annealing is below 350° C. On thecontrary, for annealing above 450° C., degradation becomes irreversible.Poly-Si-based passivating contacts are thermally more stable than SHJ asthey are fabricated at elevated temperatures (800-1000° C.) and forrather long time (>10 min). As result of the thermal treatment a buriedjunction (typical depth <200 nm) at c-Si surface with a moderate surfacedoping concentration (in the order of 10¹⁹ cm−3) is formed. The role ofthis doped region is to increase the density of the selectedcarrier-type in front the SiOx, enabling: (i) high tunneling probabilitythrough the SiOx, resulting from the better alignment of the valenceband energy positions (e.g. hole selective contact) in the c-Si and inthe heavily-doped poly-Si layer, and (ii) an efficient surfacepassivation by field effect, as the density of minority carriers (e.g.electrons for a hole selective contact) and consequently theirShockley-Read-Hall (SRH) recombination rate is reduced withoutexcessively increasing Auger recombination. Because of the lattermechanism, poly-Si based passivating contacts exhibit junctionproperties more resilient against variation of the Dit.

As a second requirement, the FlaSH contact has to be able to deliverhigh surface passivation and charge carrier transport upon theapplication of firing. The combination of thermal treatment and layerstack configuration (ie. SiOx/nc-SiC or SiOx/nc-Si(i)/nc-SiC) aredesigned to prevent the formation of doping in-diffusion into the Siwafer. Therefore, full-area contact functionalities are provided by: (1)high chemical passivation of the c-Si/SiOx interface, analogously as fora-Si:H heterojunction, and (2) a strong field effect enabling to drivethe c-Si surface in accumulation conditions. As consequence, a regionwith high density of majority carriers is induced in front the thinSiOx, which similarly to the buried junction of a diffused poly-Si basedpassivating contact enables to reduce SHR recombination and to fosterquantum tunneling of majority carriers through the thin SiOx.

To fulfill criteria (1), we placed a thin SiOx layer in between the c-Siand the deposited SiCx(p) exploiting the functionalities of: preventingepitaxial re-growth of the deposited layer on the c-Si wafer duringfiring, an effect responsible for additional surface defects; physicallydisplacing structural defects (i.e. within the deposited SiCx(p) and/orat the SiCx(p)/metal interface) from the c-Si surface; profiting of thenature of the Si-SiOx interface, which can attain very low defectdensities after high temperature thermal treatment andpost-hydrogenation. To satisfy criteria (2) we developed high workfunction nano-crystalline silicon carbide (nc-SiCx(p)) as hole selectivecontact. Such layer was obtained by depositing highly-doped a-SiCx(p):Hand fostering its crystallization into a nc-SiCx(p), during the firingprocess. The resulting highly-crystalline film can, in fact, accommodatea larger number of active dopants and thus exhibit a higher workfunction compared to only amorphous layers.

To induce surface band bending in a similar way as provided byin-diffused region, fixed charges trapped at SiOx/Si interface can alsobe used.

Surface passivation of the FlaSH contact technology was investigated onsymmetrical test structures as sketched in FIG. 30. Recombinationparameter such as effective minority carrier lifetime (τeff), impliedopen circuit voltage (i-VOC) and dark saturation current density (J0)were measured with photo-conductance decay method (PCD). The curvesdepict τeff at an injection level of 10¹⁵ cm−3 (τeff_(@)10¹⁵ cm⁻³), asfunction of the normalized CH4 over total gas flow ratio used fordepositing of the a-SiCx(p):H layers. The figure reveals an increase ofτeff with decreasing of the normalized flow ratio, until reaching anoptimum value τeff_(@)10¹⁵ cm⁻³=1100 μs beyond which it decreases again.The poor surface passivation for high flow ratio SiCx(p) layer (i.e.high C-content) might be attributed to several reasons: (i) C releasedfrom the SiCx-layer during firing reacts with oxygen from the SiOxlayer, reducing the ability of the SiOx layer to passivate the surface.An increase of the defect density and a decrease of the dopingefficiency in amorphous and nano-crystalline layers are expected at highC-content. The former effect is responsible for additional recombinationof minority charge carriers in the defective SiCx layer, or trapping offree charge carriers by the defect states in the SiCx layers, whichmight further reduce effective doping in C-rich SiCx(p) and thusadditional decrease of the work function. As result, recombination rateincreases as surface band-bending and thus repulsion of minoritycarriers from the Si surface becomes less effective. For the normalizedflow ratio of 0.18 we achieved excellent τeff_(@)10¹⁵ cm−3 above 1100 μsand i-VOC of 707 mV corresponding to J0 of 15 fA/cm2. To suppresseventual detrimental effects resulting from the interaction between Cand B atoms of the SiCx(p) layer with the chemical SiOx layer, wereplaced the first 10 nm of the SiCx(p) layer by an a-Si(i):H bufferlayer, as sketched in FIG. 30.

As FIG. 30 indicates, the a-Si(i) buffer layer efficiently fosterssurface passivation, boosting the i-VOC voltage for all conditions.Interestingly, the impact of this layer on the passivation quality isonly minor for samples with high C-content, whilst it is moresignificant for samples with low C-content. We speculate that thebenefits of the intrinsic buffer layer become evident when a sufficientband-bending at c-Si interface is provided by the nc-SiCx(p). For theFlaSH contact based on the nc-Si(i)/nc-SiCx(p) stack with normalizedflow ratio of 0.18 (for the SiCx(p)) we achieved τeff@ 10¹⁵ cm−3>1300 μsand i-VOC of 715 mV and a J0 of ˜12 fA/cm2. Without carbon alloying(i.e. normalized flow ratio equal to zero) the entire layer stackblistered during the fabrication sequence, for both with and withoutbuffer layer.

This illustrates that C incorporation within the Si network is anessential feature of the FlaSH contact as it allows obtaining layersthat are mechanically stable to firing. To conclude, FlaSH contactsbased on SiCx(p) with and without intrinsic buffer layer showed J0 equalor below 15 fA/cm2, compared to 35 fA/cm2 reached for world record PERCcells. The transport properties of the FlaSH contact were evaluated bymeans of contact resistivity (ρc) measured via transfer length method(TLM). The measured ρc strongly increased towards higher C-content forboth nc-SiCx(p) and nc-Si(i)/nc-SiCx(p) stack, as reported 30.Interestingly, for low C-content, the additional nc-Si(i) buffer layerdoes not have a strong impact on ρc. This result can be explained by thefact that the nc-SiCx(p) induces a band-banding which effectivelyrenders the nc-Si(i) buffer layer p-doped. Reasonably low ρc of 65mΩ·cm2 and 73 mΩ·cm2 for the nc-SiCx(p) and nc-Si(i)/nc-SiCx(p) stacks,(with normalized flow ratio of 0.18 for the nc-SiCx(p)) deposited on theSiOx, were respectively achieved.

The passivation and transport results here reported, demonstrate thatthe FlaSH contact is capable to deliver high carrier selectivity whilebeing fabricated by using firing thermal process.

As shown in FIG. 31, the buffer layer drastically decreases the carrierrecombination rate of the contact. Excellent injection behavior of thecontact layer is demonstrated by the high implied fill factor (i-FF).

FIG. 32A to 32F represent secondary-ion mass spectrometry (SIMS)measurements of the O, B and C doping profiles expressed in intensity[counts/seconds] as function of the depth. The SIMS depth profiles wereobtained using a CAMECA SC-Ultra instrument with O2+ primary ions with 1keV impact energy. The primary ion current was set to 2 nA. The ion beamwas scanned over an area of 250 μm×250 μm and the signal was acquiredfrom inner area 60 μm in diameter. The acceptance energy window was 19eV. The mass resolving power was set to 1200 in order to eliminate thecontribution of 10B1H+ on 11B+ signal. The secondary ions collectedwere: 11B+, 12C+, 16O+ and 28Si+(not shown here). All the depth profileswere acquired with the same instrumental settings and analyticalconditions such that the profiles can be compared directly. The depthsof the SIMS sputtering craters were measured post-acquisition with aKLA-Tencor P17 profilometer. The sputtering time was converted to depthin nanometers assuming a linear erosion rate through the differentlayers.

FIGS. 32A to 32C firstly show, for the as-deposited and fired states,that the intensity of the B signal in the c-Si, i.e. after the oxygenpeak defining the SiOx/c-Si interface, decays of more than one order ofmagnitude within 5 nm depth. The firing step mostly influences the Bsignal into the SiOx compared to the as-deposited state. In theas-deposited state, a B signal within the first 5 nm of the c-Si waferis also observed. Because no doping diffusion is expected for theas-deposited state (processing temperature did not exceed 200° C.), theB signal into the c-Si for the as-deposited state is due to the limiteddepth resolution of the SIMS measurement. These measurements clearlyhighlight one of the novelty of the invention, which is the combinationof passivating and conductive layer stack and the thermal treatmentpreventing the formation of an in-diffused doped region atlight-absorbing silicon wafer surface.

Additionally, as shown in FIG. 32D to 32F an intrinsic Si-based bufferlayer, lowers the intensity of the B signal in the SiOx layer which isbeneficial to reduce surface defect density at SiOx/c-Si interface andin an enhanced passivation as shown in FIG. 26 and by the cell resultsin FIG. 37C. The intensity of the B and O peaks at c-Si/SiOx interfaceare reported in the tables on the right side of FIG. 32.

Further investigations were also carried out by the inventors atmicrostructural level to understand the phenomena occurring in the newFlaSH contact structure of the present invention.

The microstructure of the single nanocrystalline nc-SiCx(p) and bi-layernc-Si(i)/nc-SiC(p) stacks forming the FlaSH contact, was investigatedafter firing by high resolution transmission electron microscopy (TEM).In all cases, the normalized flow ratio of the SiCx layers was set to0.18. Scanning TEM images, which were acquired using a high-angleannular dark-field (HAADF) detector, are shown alongsideenergy-dispersive X-ray spectroscopy (EDX) maps for the two contactarchitectures in FIGS. 33A and 33B, wherein the arrowhead marks theposition of the interface between the intrinsic and doped regions. Thesemicrographs and chemical maps highlight clearly the 1 to 2 nm (inprojection) passivating SiOx layer at the c-Si surface. In the case ofthe bi-layer stack, the evolution of the Si K, O K and C K edge EDXintensities along the vertical axis enable to distinguish the start ofthe doped region about 10 nm from the c-Si surface (arrowhead in FIG.33B).

High-resolution TEM (HRTEM) images demonstrate a difference incrystallinity between the two contact architectures (FIGS. 33C and 33D).The crystallinity of the nc-SiCx(p) single layer stack appearshomogenous, with no significant difference in the number of Sireflections observed in the Fourier transforms computed from the firstpart of the layer (the first 7 nm near the SiOx buffer layer, inset i)or the rest of the layer (inset ii). On the other hand, Fouriertransforms computed at the position of the intrinsic layer (the first 7nm near the SiOx buffer layer, inset iii) shows more Si reflections thanthe Fourier transform taken at the position of the nc-SiC(p) layer(inset iv). The presence of C hence appear to inhibit thecrystallization of Si. In the case of the double layernc-Si(i)/nc-SiC(p) stack (FIG. 33D), the presence of the underlyingintrinsic layer may slightly increase the crystalline fraction in theupper boron-doped layer by promoting epitaxial growth of Si crystalsacross the interface during firing. Indeed, an inverse Fourier transformof the Si<111> reflection highlighted in FIG. 33D shows that largecrystals spanning all the way across the two layers may form duringfiring, hence alleviating to some extent the negative influence of theintrinsic layer on carrier transport across the contact. Overall, onlySi reflections were indexed in the Fourier transforms of both contactstructures. The small amount of C is not sufficient to result in theformation of silicon carbide crystals (on the order of 2 at %, estimatedusing the Cliff-Lorimer method [Cliff, G. & Lorimer, G. W. Thequantitative analysis of thin specimens. J. Microsc. 103, 203-207(1975)], a value influenced by C surface contamination in the microscopechamber). It should also be mentioned that the use of shiny-etchedwafers, which exhibit a small roughness, results in a sight blurring ofthe interfaces as these are observed in projection.

FIG. 34 shows FTIR measurements of the SiCx(p) with flow ratio of 0.18before firing (as-deposited), after firing and after firing followed bya hydrogenation treatment. For these three samples, the FTIRmeasurements depicted in FIG. 34 reveal the presence of various peaks.The first one is centered at 630 cm−1 and is associated to Si—H waggingor rocking modes followed by the one centered at 780 cm−1 reveals Si—Cstretching mode while at Si—O stretching modes appear at 1050 cm−1,Si-Hn (with n=1, 2, 3) stretching modes appear in the band centered at2050 cm−1. Finally, the absorption band between 2800 and 3000 cm−1 isassociated to stretching modes of C-Hn (n=1, 2, 3) groups in sp2 and sp3orbital configurations. For the as-deposited sample, the intensity ofthe C-Hn signal is weaker compared to that of the Si—H and Si-Hn bonds.This indicates that most of the H incorporated in the a-SiCx(p):H, isbonded to Si rather than C. As expected after firing, H weakly bonded toSi effused out from the sample. As shown by the FTIR measurementreported FIG. 34 for the fired nc-SiCx(p) the Si-Hn modes are absent.Interestingly, the more temperature stable C-Hn stretching modes exhibita reduced intensity, compared to the as-deposited state but are stillpresent.

Additionally, for the fired nc-SiCx(p) a strong increase of the Si—Cpeak is observed compared to the as-deposited a-SiCx(p), suggesting theadditional Si—C bonds are built upon the thermal treatment. The shift ofSi—C vibrational mode towards higher wavenumber is a clear signature forthe occurrence of the phase transition from amorphous to crystallinestate of the film due to firing. The structural properties of thechemical SiOx also seems to be influenced by the firing process as theSi—O peak moves towards higher wavenumber indicating the formation of amore stoichiometric layer. The effect of the hydrogenation, shown inFIG. 34 appears as an increased intensity for the C-Hn band returning onthe value assumed in the as-deposited state. Eventual Si-Hn bondsresulting from the hydrogenation process are not visible, most luckilybecause below the sensitivity of the measurement system.

FIG. 39 reports FTIR measurements for SiCx(B) samples prepared withdifferent CH₄ flows, with and without buffer layer after firing andhydrogenation. The carbon content increases monotonously with the CH₄flow in the gas mix, higher CH4 flow resulting in higher carbonconcentration. The figure clearly shows that the intensity of the Si—Cpeak and C—H_(n) groups increase when going to higher CH₄ flows. To thisphenomena is also associated a reduction of the passivation quality. Thepresence of the buffer layer leads to a slight reduction of theintensity of the Si—C peak and C—H_(n) groups compared to the case whereno buffer is present.

For a deeper understanding of the H-evolution during the fabricationsequence of the FlaSH contact we have performed TDS measurements. Asreported in literature H-effusion spectra are strongly depending on themicrostructural properties of the a-Si:H layer. Typically, two effusionprocesses can be observed, one at high temperature (HT) (i.e. >500° C.),and one at low temperature (LT) (<400° C.) [Beyer, W. Diffusion andevolution of hydrogen in hydrogenated amorphous and microcrystallinesilicon. Sol. Energy Mater. Sol. Cells 78, 235-267 (2003)]. The HT peakis attributed to diffusion of mono-hydride (atomic hydrogen) through thelayer network. The LT peak is associated to H2+ desorption due tosimultaneous rupture of two Si—H bonds forming di-hydride (molecularhydrogen) that effuses out through internal voids or layer rupture.Diffusion of atomic H is the dominant mechanism of hydrogen transport inmaterials with a dense network, thus limiting effusion rate. On thecontrary, for materials with a microstructure formed of interconnectedvoids, H2 effusion through the network becomes dominant. Effusionspectra showing the presence of both phenomena are explained by areconfiguration of the internal Si—Si bonds, occurring during theheating process, transforming the network of voids in a more compactone. The position and intensity of both HT and LT peaks can be furtherinfluenced by the composition of the a-Si alloy. For instance, theintensity of the LT peak is reported to increase with adding carbon tothe Si network due the formation of a higher density of voids. Inaddition, LT and HT effusion peaks are Fermi-level dependent as itsenergy position influences the rapture energy of the Si—H bonds. Forthis reason, a-Si:H(p) show a strong shift of both effusion peakstowards lower temperature compared to intrinsic a-Si(i):H. Massseparated ions (m/z) m/z=2 associated to H2 effusion for the FlaSHcontact composed by the SiCx(p) with a normalized flow ratio of 0.18were measured in the as-deposited, fired and after hydrogenation.

As shown in FIG. 35, the a-SiCx(p):H (in the as-deposited state) shows abroad H2+ effusion spectra indicating the presence of both effusionpeaks with the LT at 275° C. and HT at 550° C. which appears as ashoulder or plateau. The presence of the LT peaks is a signature of H2desorption through interconnected voids within the a-SiCx(p):H layer.The decrease of the effusion rate towards higher temperatures isassociated to the collapse of the void network in favor of a morecompact one where diffusion of atomic H becomes dominant, thus limitingthe effusion rate. Upon firing at 800° C., H effusion form the nc-SiC(p)is, as expected, much lower compared to the as-deposited a-SiCx(p):H(FIG. 35). An increased desorption is observed towards highertemperatures and might be associated to the rupture of the moretemperature-stable C—H bonds (see FTIR measurements in FIG. 35). Them/z=2 spectra reported for the hydrogenated nc-SiCx(p) sample, shows ahigher H-effusion with respect to after firing proving that H is indeedreintroduced. Additionally, the effusion signal exhibits a weak LT peakat ˜350° C. and a slight increase towards higher temperatures. Such LTeffusion process is shifted towards higher temperature compared to theone in the as-deposited state. This effect might be explained by twophenomena: hydrogen is mainly reintroduced at SiOx/Si interface andpossibly forms O—H bonds that are more thermally stable than Si—H;and/or the more compact network of the nc-SiCx(p) compared to the asdeposited a-SiCx(p):H, limiting H2 effusion rate.

The effect of light soaking was also investigated. A remarkablecharacter for symmetrical samples is that they show an improvement ofthe surface passivation quality after light exposure.

Three samples with SiCx(p) layers with flow ratio of 0.78 and 0.18 withand without Si(i) as buffer layer were light soaked at 1 sunillumination. After light soaking the samples were stored in dark for 1month. As FIG. 36 clearly indicates, τ_(eff) and i-Voc stronglyincreases after the first 2 seconds of light soaking. For longerexposure time no substantial changes in the surface passivation isobserved. After storing the samples in the dark for one month, τ_(eff)and i-Voc were mostly unchanged.

Therefore, the light soaking effect is most likely to affect the densityof charges at SiOx/Si during illumination. In more detail, danglingbonds at SiOx/Si interface lead to trap states. Their energydistribution consists of two different peaks located within the Si bandgap. In particular, traps energetically located above the energy of theintrinsic Fermi in Si (E_(fi)) level can act as donor state, and areneutral when empty, and negatively charged when filled by an electron.On the contrary, trap states located at energies below E_(fi), act asacceptor states and positive when empty, and neutral when filled by anelectron. This makes the charge state of the interfacial traps to bedependent on the Fermi level position. For the FlaSH contact of theinvention, empty donor states and H2+ at Si/SiOx interface areresponsible for positive charges. Such positive fixed charges attractelectrons towards the interface increasing the Shockley-red-hallrecombination rate. Electrons injected during illumination can fillempty donor states of deactivate H2+ reducing the density of fixedcharges which results in a beneficial band bending.

This light soaking effect also indicates that the band-bending providedby the work function of the SiCx(p) and fixed charges is stronger thanthe one provided by the ultra-shallow doping in-diffusion region. Thechange of surface band bending before and after light soaking was alsomeasured by surface photo voltage measurements as reported in FIG. 42.From such measurements can be derived that after light soaking, thedensity of positive fixed charges decreases.

Hybrid solar cells demonstrators were built to show the workingpotential of the FlaSH heterojunction contact of the present invention,including hybrid cells involving single stack SiCx(B) (FIG. 37A) ora-Si(i)/SiCx(B) (FIG. 37B) as hole selective contact on the rear sideand involving a-Si(i):H/a-Si(n):H as front side emitter. FIG. 38Crepresents J-V measurements of the hybrid cells involving single stackSiCx(B) or a-Si(i)/SiCx(B) as hole selective fired contact on the rearside. Table 2 below shows a summary of the external parameters of thehybrid cells in each case:

TABLE 2 J_(SC) V_(OC) FF η p-FF [mA/cm²] [mV] [%] [%] [%] nc-SiC_(x)(p)38.3 697 79.6 21.3 81.6 nc-Si(i)/nc-SiC_(x)(p) 38.2 705 78.5 21.1 81.8

Light soaking of the hybrid cells was conducted at 1-sun illumination byconsecutive measurements with sun simulator. As shown in FIG. 38 A-D,for solar cells a strong increase of the Voc is observed after 6consecutive measurements under AM 1.5 conditions. This result is in linewith the results already obtained on symmetrical samples, see FIG. 36.In addition to the Voc increase the cells also showed an increase of theFF. As previously discussed, it is considered that the density ofpositive charges at Si/SiOx is reduced after illumination. This effect,leads to an increases of the potential barrier for the electrons andthus reducing their recombination rate and at the same time decreasesthe potential barrier for the holes thus better transport.

The effect of the peak firing temperature (i.e. 800 and 850° C.) for thesingle SiCx(p) with normalized flow ration of 0.18 and constant firingtime of 3 sec was evaluated on hybrid cells and is presented in FIG. 43.The final conversion efficiency was found to be quite stable with thepeak firing temperature. In FIG. 44 we report the impact of thesputtering damage evaluated by PLI, the device fabricated at lowertemperature showed a lower sputtering damage before and after curing ofthe ITO (i,e, annealing at 210° C. for 30 min).

The effect of the layer doping expressed by the normalized to the maxvalue of the TMB_(2%) for the single SiCx(p) with normalized flow rationof 0.18 fired at 800° C. for 3 sec was evaluated on hybrid cells and ispresented in FIG. 45. The final conversion efficiency was found toincrease with the normalized TMB_(2%). In FIG. 46 we report the impactof the sputtering damage evaluated by PLI, the device with higherTMB_(2%) showed a lower sputtering damage before and after curing.

The effect of the peak firing time (i.e. 3 and 6 sec) for the singleSiCx(p) with normalized flow ration of 0.18 and constant firing peaktemperature of 800° C. was evaluated on hybrid cells and is presented inFIG. 47. The final conversion efficiency was found to be quite stablewith the peak firing time. In FIG. 48 we report the impact of thesputtering damage evaluated by PLI, both devices showed similarsputtering damage before and after curing of the ITO (i,e, annealing at210° C. for 30 min).

The effect of the thickness of the intrinsic buffer layer expressed bythe normalized to the max value of the deposition time for the bilayernc-Si(i)/nc-SiCx(p) with normalized flow ration of 0.18 fired at 800° C.for 3 sec was evaluated on hybrid cells and is presented in FIG. 49. Thefinal conversion efficiency was found to increase towards thicker bufferlayer. In FIG. 46 we report the impact of the sputtering damageevaluated by PLI, the device with intrinsic buffer layer showed a lowersputtering damage before and after curing. In conclusion doping, firingprocess and buffer layer also play a role on the sputtering damage ofthe heat resistant-heterojunction. In FIG. 53 alternative processfabrication of the heat resistant heterojunction including hydrogenationduring firing from a capping layer is shown. The results of the singleand bilayer stack with bilayer with normalized flow ratio of the SiCx(p)of 0.18 of the fired heat resistant heterojunction fabricated accordingto the process sketched in FIGS. 29 and 52, respectively are shown inTable 3. Similar i-VOC could be reached for both layers and bothprocesses, however, having the hydrogenation included during the firingprocess has the advantage of further reducing process complexity.Additionally the C-content of one of the layers composing the heatresistant heterojunction can be optimized in order to prevent Ag-fritpenetration during firing.

TABLE 3 Processed according to: Layer stack i-V_(OC) [mV] FIG. 29nc-SiCx(p) 705 FIG. 29 nc-Si(i)/nc-SiCx(p) 715 FIG. 51 nc-SiCx(p) 700FIG. 51 nc-Si(i)/nc-SiCx(p) 716

FlaSH Cell Integration

Industrial p-PERC solar cells are typically fabricated according to theprocess reported in top panel of FIG. 52. After texturing and cleaningof the wafer, an emitter is fabricated by phosphorous diffusion followedby SiNx:H deposition on serving as surface passivation andantireflection. A Si etch-back of the rear side including also aplanarization of the rear surface is then performed. For simplicity,this step is not reported in the fabrication sequence. The structure ofthe resulting cell precursor is depicted in step 1 of FIG. 52. Next, adielectric layer (or stack of dielectrics) is deposited on the rearside, see steps 2, and subsequently patterned as sketched in step 3.After screen-printing of Ag and Al pastes, respectively on the front andrear side of the cell (steps 4 and 5), a firing step is applied (step6). Upon firing, the front Ag-based paste sinters through the SiNx:H andestablishes electrical contacts with the underlying emitter, while onthe rear side local Al-doped Si regions are formed (step 6). As alreadydiscussed in the FlaSH represents an evolutionary upgrade of the rearside contact of a PERC cell as it can be integrated with only fewchanges of the existing process flow. As depicted in the bottom panel ofFIG. 52, the process flow of an industrial cell architecture employingthe FlaSH contact on the rear side, employs the same cell precursor asfor standard p-PERC (see step 1). The growth of the thin SiOx can beperformed during a cleaning step in hot HNO3 performed after rear sideplanarization. Afterwards, the SiCx(p)-based forming the FlaSH contactis deposited via PECVD on the rear side (step 2.a). Because the FlaSHcontact combines the functionalities of passivation and charge carriertransports, patterning of the rear contact is unnecessary. An Ag-basedgrid is then screen-printed on the front side (step 3.a), and the solarcell is then fired. As result an electrical contacts between the Agpaste and the front emitter is established while on the rear sidecrystallization of the FlaSH contact takes (a-SiCx(p):H partiallycrystalizes into a nc-SiCx(p)) as depicted in step 4.a. Hydrogenation isthen carried out in order to enhance surface passivation (step 5.a) andalternatively can be performed during the firing process. Finally, therear side could be metallized by screen-printing of Al paste or with alow cost TCO/metal stack (step 6.a).

The investigation of the optimal deposition conditions of thehole-selective contact was done on symmetrical test structuresfabricated according to the process flow presented in FIG. 29. Borondoped float zone (FZ) 4-inch silicon wafers shiny etched, (100)oriented, with thickness of 250 μm and resistivity ˜2 Ω·cm were used.After cleaning using standard wet chemistry, a thin SiO_(x) layer (˜1.2nm) was grown by wet chemical oxidation in hot nitric acid also referredto as “chemical SiO_(x)”. Afterwards, a-SiC_(x)(B):H or bilayera-Si(i):H/a-SiC_(x)(B):H were deposited via plasma enhanced chemicalvapor deposition (PECVD) on both sides of the wafer. A rapid thermalannealing process (RTA) emulating the firing process of industrial c-Sisolar cells was applied. After the firing process, all samples werecoated on both side with a hydrogen donor layer deposited via PECVD. Theeffective minority carrier lifetime was measured by photo-conductancedecay (PCD) and the method of Kane and Swanson was applied to extractthe emitter saturation current density at an excess carrier density of10¹⁶ cm⁻³. The spatial homogeneity of the passivation was analyzed usingphotoluminescence imaging (PLI). For electrical characterization,coplanar ITO/Ag contacts were sputtered, and the specific contactresistivity was measured using TLM.

Hybrid c-Si solar cells employing conventional silicon heterojunction onthe front side and the fired hole selective contact on the rear sidewere fabricated. The process flow of the fabricated solar cell is shownin FIG. 29. For the fabricated device we used FZ p-type wafers withresistivity ˜2 Ω·cm. Single side surface texturing was performed byusing a sacrificial SiNx layer on the rear side. As the figure clearlydescribes, after cleaning and chemical oxidation, the a-SiC_(x)(B):H orthe bilayer a-Si:H/a-SiC_(x)(B):H where deposited via PECVD deposition.Solar cells were then fired at T=800° C. for 3 seconds, andhydrogenation as described in the previous section was employed. Afterstripping of the hydrogen donor layer in wet chemical solution, waferswere cleaned. Front side heterojunction formed by a-Si(i):H/a-Si(n):Hwas deposited on the front side by means of PECVD. Transparentconductive oxide (ITO) was sputtered on both sides of the wafers inorder to enhance lateral conductivity and optics (i.e. antireflectioncoating and internal rear reflectance) of the solar cells. Finally, Agwas sputtered on the full rear side, while Ag metal fingers were screenprinted on the front side. Solar cell I-V measurements were performed byusing a Wacom solar simulator calibrated with reference cells measuredat Fraunhofer ISE.

1-22. (canceled)
 23. A heterojunction solar cell comprising a firstelectrode layer, a second electrode layer comprising a metallic contactlayer, a light-absorbing silicon layer and a passivating and conductivestack arranged between said light-absorbing silicon layer and at leastone of the first and second electrode layers, wherein said passivatingand conductive stack comprises a wide band gap material layer having anelectronic band gap greater than 1.4 eV and at least one dopedsilicon-based layer, said wide band gap material layer being applied ona surface of the light-absorbing silicon layer between saidlight-absorbing silicon layer and said doped silicon-based layer to forma a heat-resistant heterojunction, said heat resistant heterojunctionbeing arranged for at least maintaining its passivating and conductiveproperties after thermal treatment thereof above 600° C., and wherein aninterface between said passivating and conductive stack and the lightabsorbing silicon layer has a very abrupt distribution of doping profilein which the intensity of dopants in said light absorbing silicon layerdecays over at least one order of magnitude within a distance of lessthan 5 nm from said interface, said passivating and conductive stack andthermal treatment being configured to prevent dopants from diffusingfrom the passivating and conductive stack to the light absorbing siliconlayer.
 24. The solar cell according to claim 23, wherein the wide bandgap material layer has a thickness of at most 20 nm, preferably between0.5 and 2 nm.
 25. The solar cell according to claim 23, wherein saidwide band gap material layer comprises any of the following materials:SiO_(x), SiC_(x), AlO_(x), HfO_(x), AlHfO_(x), AlN, TiN, SiN_(x). 26.The solar cell according to claim 23, wherein said doped silicon-basedlayer has an atomic hydrogen concentration of less than 5%, defined asthe number of hydrogen atoms per unit volume divided by the total numberof all atoms per unit volume of the doped silicon-based layer.
 27. Thesolar cell according to claim 23, wherein said doped silicon-based layeris a doped-silicon carbide based layer SiC_(x).
 28. The solar cellaccording to claim 24, characterized in that it further comprises atleast one buffer layer arranged between said wide band gap materiallayer and said doped silicon-based layer, said buffer layer beingarranged to tune the density of dopants diffusing from the silicon-dopedlayer into the wide band gap material layer and light absorbing siliconlayer.
 29. The solar cell according to claim 28, wherein said bufferlayer is made of at least one of the materials chosen among a siliconlayer Si, SiC_(x), SiN_(x), SiO_(x), SiC_(x)N_(y), SiC_(x)O_(y),SiN_(x)O_(y), SiC_(x)N_(y)Oz or a combination thereof.
 30. The solarcell according to claim 23, wherein a transparent conductive oxide layeris arranged between said doped silicon layer and said metallic layer.31. The solar cell according to claim 23, wherein the surface of thelight absorbing silicon layer whereupon the wide band gap material layeris applied is structured.
 32. The solar cell according to claim 31,wherein said wide band gap material layer comprises through-holesextending from a first surface to a second surface.
 33. The solar cellaccording to claim 23, wherein a capping layer is arranged between saiddoped silicon-based layer and said metallic contact layer.
 34. The solarcell according claim 23, wherein the first electrode layer is a frontlayer and the second electrode layer is a back layer, said lightabsorbing silicon layer, said doped silicon-based layer and said wideband gap layer being arranged between said first and second electrodelayers.
 35. The solar cell according to claim 23, wherein the first andthe second electrode layers are arranged on a same one side of saidlight absorbing silicon layer.
 36. The solar cell according to claim 33,wherein the passivating and conductive stack and/or the capping layercomprise(s) passivating species to passivate defects within the saidpassivating and conductive stack and/or at the interface between saidpassivating and conductive stack and said light absorbing silicon layer,said passivating species being releasable upon said thermal treatment.37. A heterojunction solar cell comprising a first electrode layer, asecond electrode layer comprising a metallic contact layer, alight-absorbing silicon layer and a passivating and conductive stackarranged between said light-absorbing silicon layer and at least one ofthe first and second electrode layers, wherein said passivating andconductive stack comprises a wide band gap material layer having anelectronic band gap greater than 1.4 eV and at least one dopedsilicon-based layer, said wide band gap material layer being applied ona surface of the light-absorbing silicon layer between saidlight-absorbing silicon layer and said doped silicon-based layer to forma a heat-resistant heterojunction, said heat resistant heterojunctionbeing arranged for at least maintaining its passivating and conductiveproperties after thermal treatment thereof above 600° C., wherein thelight absorbing silicon layer comprises a doped region with sheetresistance between 1 and 10⁵ ohm/square at an interface between saidpassivating and conductive stack and the light absorbing silicon layer,said doped region being obtained by diffusion of dopants from any of thelayers forming the passivating and conductive stack into to the lightabsorbing silicon layer during said thermal treatment, wherein the wideband gap material layer has a thickness of at most 20 nm, preferablybetween 0.5 and 2 nm, wherein said wide band gap material layercomprises any of the following materials: SiO_(x), SiC_(x), AlO_(x),HfO_(x), AlHfO_(x), AlN, TiN, SiN_(x), wherein said doped silicon-basedlayer has an atomic hydrogen concentration of less than 5%, defined asthe number of hydrogen atoms per unit volume divided by the total numberof all atoms per unit volume of the doped silicon-based layer, whereinsaid heterojunction solar cell further comprises at least one bufferlayerarranged between said wide band gap material layer and said dopedsilicon-based layer, said buffer layer being arranged to tune thedensity of dopants diffusing from the silicon-doped layer into the wideband gap material layer and light absorbing silicon layer, wherein saidbuffer layer is made of at least one of the materials chosen among asilicon layer Si, SiC_(x), SiN_(x), SiO_(x), SiC_(x)N_(y), SiC_(x)O_(y),SiN_(x)O_(y), SiC_(x)N_(y)O_(z) or a combination thereof, wherein saiddoped silicon-based layer is a doped-silicon carbide based layerSiC_(x), and wherein said doped silicon-based layer comprises fluorine.38. A method for manufacturing a solar cell according to claim 23,comprising the steps of: Providing a light absorbing silicon layer, andForming a passivating and conductive stack on a surface of the lightabsorbing silicon layer, said passivating and conductive stackcomprising a wide band gap material layer deposited on a surface of thelight absorbing silicon layer and a doped silicon based layer on saidwide band gap material layer, arranged between said doped silicon-basedlayer and said metallic contact layer and, thermally treating thepassivating and conductive stack at above 600° C., in order to releasespassivating species from the said passivating and conductive stackand/or at an interface between said passivation stack and said lightabsorbing silicon layer.
 39. A method according to claim 38, wherein thethermal treatment comprises a firing step at temperatures above 600° C.and with at temperature ramp rates of more than 20° C./s and dwell timeat maximum temperature below 10 s.
 40. A method according to claim 38,wherein the wide band gap material layer is made of SiO_(x) depositedvia chemical, plasma, gas-phase or light-excited methods, with x beingchosen between 0.5 and
 2. 41. A method according to claim 38, whereinfurther comprising the steps of: forming a capping layer between saiddoped silicon-based layer of the passivating and conductive stack andsaid metallic contact layer; and thermally treating the passivating andconductive stack and the capping layer in order to release passivatingspecies from any of the passivating and conductive stack and the cappinglayer to passivate defects within the said passivating and conductivestack and/or at the interface between said passivating and conductivestack and said light absorbing silicon layer.
 42. A method formanufacturing a solar cell according to claim 37, comprising the stepsof: Providing a light absorbing silicon layer, and Forming a passivatingand conductive stack on a surface of the light absorbing silicon layer,said passivating and conductive stack comprising a wide band gapmaterial layer deposited on a surface of the light absorbing siliconlayer and a doped silicon based layer on said wide band gap materiallayer, arranged between said doped silicon-based layer and said metalliccontact layer and, thermally treating the passivating and conductivestack at above 600° C., in order to releases passivating species fromthe said passivating and conductive stack and/or at an interface betweensaid passivation stack and said light absorbing silicon layer.
 43. Amethod according to claim 42, wherein the thermal treatment comprises afiring step at temperatures above 600° C. and with at temperature ramprates of more than 20° C./s and dwell time at maximum temperature below10 s.
 44. A method according to claim 42, wherein the wide band gapmaterial layer is made of SiO_(x) deposited via chemical, plasma,gas-phase or light-excited methods, with x being chosen between 0.5 and2.
 45. A method according to claim 42, wherein further comprising thesteps of: forming a capping layer between said doped silicon-based layerof the passivating and conductive stack and said metallic contact layer;and thermally treating the passivating and conductive stack and thecapping layer in order to release passivating species from any of thepassivating and conductive stack and the capping layer to passivatedefects within the said passivating and conductive stack and/or at theinterface between said passivating and conductive stack and said lightabsorbing silicon layer.